Metal Matrix Composites 3038970875, 9783038970873

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Metal Matrix Composites Edited by

Manoj Gupta Printed Edition of the Special Issue Published in Metals

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Metal Matrix Composites

Metal Matrix Composites

Special Issue Editor Manoj Gupta

MDPI • Basel • Beijing • Wuhan • Barcelona • Belgrade

Special Issue Editor Manoj Gupta Materials Group, Department of Mechanical Engineering Singapore Editorial Office MDPI St. Alban-Anlage 66 Basel, Switzerland

This edition is a reprint of the Special Issue published online in the open access journal Metals (ISSN 2075-4701) from 2016–2018 (available at: http://www.mdpi.com/journal/metals/special_issues/matrix_composites). For citation purposes, cite each article independently as indicated on the article page online and as indicated below:

LastName, A.A.; LastName, B.B.; LastName, C.C. Article title. Journal Name Year. Article number, Page Range.

ISBN 978-3-03897-087-3 (Pbk) ISBN 978-3-03897-088-0 (PDF)

Cover photo courtesy of Khin Sandar Tun.

Articles in this volume are Open Access and distributed under the Creative Commons Attribution license (CC BY), which allows users to download, copy and build upon published articles even for commercial purposes, as long as the author and publisher are properly credited, which ensures maximum dissemination and a wider impact of our publications. The book taken as a whole is © 2018 MDPI, Basel, Switzerland, distributed under the terms and conditions of the Creative Commons license CC BY-NC-ND (http://creativecommons.org/licenses/by-nc-nd/4.0/).

Contents About the Special Issue Editor . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . . .

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Manoj Gupta Metal Matrix Composites Reprinted from: Metals 2018, 8, 379, doi: 10.3390/met8060379 . . . . . . . . . . . . . . . . . . . . .

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Abolfazl Azarniya, Mir Saman Safavi, Saeed Sovizi, Amir Azarniya, Biao Chen, Hamid Reza Madaah Hosseini and Seeram Ramakrishna Metallurgical Challenges in Carbon Nanotube-Reinforced Metal Matrix Nanocomposites Reprinted from: Metals 2017, 7, 384, doi: 10.3390/met7100384 . . . . . . . . . . . . . . . . . . . . .

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Khin Sandar Tun, Yuming Zhang, Gururaj Parande, Vyasaraj Manakari and Manoj Gupta Enhancing the Hardness and Compressive Response of Magnesium Using Complex Composition Alloy Reinforcement Reprinted from: Metals 2018, 8, 276, doi: 10.3390/met8040276 . . . . . . . . . . . . . . . . . . . . . 47 Weidong Song, Liansong Dai, Lijun Xiao, Cheng Wang, Xiaonan Mao and Huiping Tang A Meso-Mechanical Constitutive Model of Particle-Reinforced Titanium Matrix Composites at High Temperatures Reprinted from: Metals 2017, 7, 15, doi: 10.3390/met7010015 . . . . . . . . . . . . . . . . . . . . . 57 Feng Qiu, Xiang Gao, Jian Tang, Yu-Yang Gao, Shi-Li Shu, Xue Han, Qiang Li and Qi-Chuan Jiang Microstructures and Tensile Properties of Al–Cu Matrix Composites Reinforced with Nano-Sized SiCp Fabricated by Semisolid Stirring Process Reprinted from: Metals 2017, 7, 49, doi: 10.3390/met7020049 . . . . . . . . . . . . . . . . . . . . . 69 Amit Kumar, Khin Sandar Tun, Amit Devendra Kohadkar and Manoj Gupta Improved Compressive, Damping and Coefficient of Thermal Expansion Response of Mg–3Al–2.5La Alloy Using Y2 O3 Nano Reinforcement Reprinted from: Metals 2017, 7, 104, doi: 10.3390/met7030104 . . . . . . . . . . . . . . . . . . . . . 77 Nguyen Thi Hoang Oanh, Nguyen Hoang Viet, Ji-Soon Kim and Alberto Moreira Jorge Junior Characterization of In-Situ Cu–TiH2–C and Cu–Ti–C Nanocomposites Produced by Mechanical Milling and Spark Plasma Sintering Reprinted from: Metals 2017, 7, 117, doi: 10.3390/met7040117 . . . . . . . . . . . . . . . . . . . . . 88 Ahmed Nassef, Waleed H. El-Garaihy and Medhat El-Hadek Characteristics of Cold and Hot Pressed Iron Aluminum Powder Metallurgical Alloys Reprinted from: Metals 2017, 7, 170, doi: 10.3390/met7050170 . . . . . . . . . . . . . . . . . . . . . 100 Ahmed Nassef, Waleed H. El-Garaihy and Medhat El-Hadek Mechanical and Corrosion Behavior of Al-Zn-Cr Family Alloys Reprinted from: Metals 2017, 7, 171, doi: 10.3390/met7050171 . . . . . . . . . . . . . . . . . . . . . 112 Hyun Min Nam, Duck Min Seo, Hyung Duk Yun, Gurunathan Thangavel, Lee Soon Park and Su Yong Nam Transparent Conducting Film Fabricated by Metal Mesh Method with Ag and Cu@Ag Mixture Nanoparticle Pastes Reprinted from: Metals 2017, 7, 176, doi: 10.3390/met7050176 . . . . . . . . . . . . . . . . . . . . . 124 v

Linhui Zhang, Yan Jiang, Qianfeng Fang, Rui Liu, Zhuoming Xie, Tao Zhang, Xianping Wang and Changsong Liu Comparative Investigation of Tungsten Fibre Nets Reinforced Tungsten Composite Fabricated by Three Different Methods Reprinted from: Metals 2017, 7, 249, doi: 10.3390/met7070249 . . . . . . . . . . . . . . . . . . . . . 132 Sonia ´ Simoes, ˜ Filomena Viana, Marcos A. L. Reis and Manuel F. Vieira Aluminum and Nickel Matrix Composites Reinforced by CNTs: Dispersion/Mixture by Ultrasonication Reprinted from: Metals 2017, 7, 279, doi: 10.3390/met7070279 . . . . . . . . . . . . . . . . . . . . . 144 Milli Suchita Kujur, Ashis Mallick, Vyasaraj Manakari, Gururaj Parande, Khin Sandar Tun and Manoj Gupta Significantly Enhancing the Ignition/Compression/Damping Response of Monolithic Magnesium by Addition of Sm2 O3 Nanoparticles Reprinted from: Metals 2017, 7, 357, doi: 10.3390/met7090357 . . . . . . . . . . . . . . . . . . . . . 155 Hajo Dieringa, Lydia Katsarou, Ricardo Buzolin, G´abor Szak´acs, Manfred Horstmann, Martin Wolff, Chamini Mendis, Sergey Vorozhtsov and David StJohn Ultrasound Assisted Casting of an AM60 Based Metal Matrix Nanocomposite, Its Properties, and Recyclability Reprinted from: Metals 2017, 7, 388, doi: 10.3390/met7100388 . . . . . . . . . . . . . . . . . . . . . 172 Cristina Ar´evalo, Isabel Montealegre-Melendez, Eva M. P´erez-Soriano, Enrique Ariza, Michael Kitzmantel and Erich Neubauer Study of the Influence of TiB Content and Temperature in the Properties of In Situ Titanium Matrix Composites Reprinted from: Metals 2017, 7, 457, doi: 10.3390/met7110457 . . . . . . . . . . . . . . . . . . . . . 185 Youhong Sun, Chi Zhang, Baochang Liu, Qingnan Meng, Shaoming Ma and Wenhao Dai Reduced Graphene Oxide Reinforced 7075 Al Matrix Composites: Powder Synthesis and Mechanical Properties Reprinted from: Metals 2017, 7, 499, doi: 10.3390/met7110499 . . . . . . . . . . . . . . . . . . . . . 198 Abdollah Saboori, Seyed Kiomars Moheimani, Matteo Pavese, Claudio Badini and Paolo Fino New Nanocomposite Materials with Improved Mechanical Strength and Tailored Coefficient of Thermal Expansion for Electro-Packaging Applications Reprinted from: Metals 2017, 7, 536, doi: 10.3390/met7120536 . . . . . . . . . . . . . . . . . . . . . 212 Suqing Zhang, Tijun Chen, Jixue Zhou, Dapeng Xiu, Tao Li and Kaiming Cheng Mechanical Properties of Thixoforged In Situ Mg2 Sip /AM60B Composite at Elevated Temperatures Reprinted from: Metals 2018, 8, 106, doi: 10.3390/met8020106 . . . . . . . . . . . . . . . . . . . . . 226 Josef Zapletal, Zuzanka Trojanov´a, Pavel Dolevolumezal, Stanislava Fintov´a and Michal Knapek Elastic and Plastic Behavior of the QE22 Magnesium Alloy Reinforced with Short Saffil Fibers and SiC Particles Reprinted from: Metals 2018, 8, 133, doi: 10.3390/met8020133 . . . . . . . . . . . . . . . . . . . . . 239 Sravya Tekumalla, Najib Farhan, Tirumalai S. Srivatsan and Manoj Gupta Nano-ZnO Particles’ Effect in Improving the Mechanical Response of Mg-3Al-0.4Ce Alloy Reprinted from: Metals 2016, 6, 276, doi: 10.3390/met6110276 . . . . . . . . . . . . . . . . . . . . . 252 vi

Gururaj Parande, Vyasaraj Manakari, Harshit Gupta and Manoj Gupta Magnesium-β-Tricalcium Phosphate Composites as a Potential Orthopedic Implant: A Mechanical/Damping/Immersion Perspective Reprinted from: Metals 2018, 8, 343, doi: 10.3390/met8050343 . . . . . . . . . . . . . . . . . . . . . 263

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About the Special Issue Editor Manoj Gupta was a former Head of the Materials Division of the Mechanical Engineering Department and Director designate of the Materials Science and Engineering Initiative at NUS, Singapore. He obtained his Ph.D. from the University of California, Irvine, USA (1992) and conducted postdoctoral research at the University of Alberta, Canada (1992). In August 2017, he was highlighted among the top 1% scientists by The Universal Scientific Education and Research Network and among the top 2.5% scientists by ResearchGate. He developed the Disintegrated Melt Deposition technique and the Hybrid Microwave Sintering technique, an energy-efficient, solid-state processing method to synthesize alloys/micro/nano-composites. He has published over 490 peer-reviewed journal papers and owns two US patents. His current h-index is 56 and his RG index is 46; his publications have received more than 13,000 citations. He has also co-authored six books, published by John Wiley, Springer, and MRF-USA. He is Editor-in-Chief and Editor of 12 international peer-reviewed journals. In 2018, he was announced World Academy Championship Winner in the area of Biomedical Sciences by the International Agency for Standards and Ratings. A multiple award winner, he actively collaborates with various scientist and visits Japan, France, Saudi Arabia, Qatar, China, USA, and India.

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metals Editorial

Metal Matrix Composites Manoj Gupta Materials Group, Department of Mechanical Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore 117576, Singapore; [email protected] Received: 21 May 2018; Accepted: 23 May 2018; Published: 24 May 2018

Metal Matrix Composites (MMCs) are a unique class of materials capable of providing design freedom to material scientists, allowing the creation of materials that can be targeted to a wide spectrum of applications [1,2]. The key factors that affect the design of MMCs include the following [1–3]: (a) (b) (c) (d)

Choice of matrix material. Type, size (length scale) and amount of reinforcement. Type of processing (primary and secondary), as this controls the microstructure including matrix-reinforcement integrity. Heat treatment procedure.

The choice of matrix material and reinforcement is primarily influenced by the end application. For example, nickel-based materials are commonly chosen as matrices for high temperature applications, while titanium- and magnesium-based materials are considered for both light weighting of engineering structures and biomedical applications [1–5]. Similarly, different processing methods (primary and secondary) are chosen depending on many factors, such as cost of the end product, type of microstructure desired and volume of production. As an example, conventional casting is always preferred when high volume production is required. The proper utilization of heat treatment also plays a crucial role in selectively enhancing the properties of composites. In view of the dynamic scientific and application potential of metal matrix composites, the present thematic issue was launched and was highly successful, with 20 papers contributed by researchers from all over the world and accepted after rigorous peer review. Key materials that were investigated include the following: (a) (b) (c) (d) (e) (f)

Aluminum-based materials. Magnesium-based materials. Titanium-based materials. Copper-based materials. Silver-based materials. Tungsten-based materials. Reinforcement investigated in the above matrices included at least one of the following:

(a) (b) (c) (d) (e) (f) (g) (h) (i)

ZnO in nanolength scale. SiC in nanolength scale. Y2 O3 in nanolength scale. Sm2 O3 in nanolength scale. Carbon-based reinforcements such as CNTs and graphene oxide. Metal-based reinforcement such as complex alloy reinforcement in micron length scale. Tungsten fiber nets. TiB. Intermetallic reinforcement such as Mg2 Si.

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Metals 2018, 8, 379

(j) (k)

Saffil fibers. β-Tricalcium phosphate. The types of processing that were employed to synthesize composites included the following:

(a) (b) (c) (d) (e)

Liquid based methods such as the disintegrated melt deposition method including ultrasonication. Semi-solid or two-phase methods. In-situ method. Solid state method including cold and hot pressing, mechanical milling and spark plasma sintering. Thixoforging. The reported properties of the metal matrix composites in this special issue included:

(a) (b) (c) (d) (e) (f)

Mechanical properties such as tensile and compression response. Ignition response. Damping response. Coefficients of thermal expansion. Elastic and plastic deformation behavior. Corrosion response.

The choice of matrix, reinforcement, processing method and characterization results reported by researchers in the twenty papers included in this special issue clearly indicate the prevalence of scientific curiosity and the optimism of researchers worldwide to create new metal-based composite materials for a wide range of applications, spanning from the biological sector to multiple engineering sectors. This special issue certainly provides a succinct description of research activities conducted across the world and thus will be very useful for students and researchers in both academia and industry. Finally, I would like to thank all the authors for their excellent contributions to this issue, to the reviewers for making useful comments, and to the Metals editorial staff for publishing these articles promptly. Conflicts of Interest: The authors declare no conflict of interest.

References 1. 2. 3.

4. 5.

Ibrahim, I.A.; Mohamed, F.A.; Lavernia, E.J. Particulate reinforced metal matrix composites—A review. J. Mater. Sci. 1991, 26, 1137–1156. [CrossRef] Lloyd, D.J. Particle reinforced aluminium and magnesium matrix composites. Int. Mater. Rev. 1994, 39, 1–23. [CrossRef] Ceschini, L.; Dahle, A.; Gupta, M.; Jarfors, A.E.W.; Jayalakshmi, S.; Morri, A.; Rotundo, F.; Toschi, S.; Singh, R.A. Aluminum and Magnesium Metal Matrix Nanocomposites; Springer: Singapore, 2016; ISBN 978-981-10-2680-5 (Print); 978-981-10-2681-2 (Online). Gupta, M.; Seetharaman, S. Magnesium Based Nanocomposites for Cleaner Transport. In Nanotechnology for Energy Sustainability; Raj, B., Van de Voorde, M., Mahajan, Y., Eds.; Wiley-VCH: Weinheim, Germany, 2017. Insight into Designing Biocompatible Magnesium Alloys and Composites; Gupta, M., Meenashisundaram, G.K., Eds.; Springer: Singapore, 2015. © 2018 by the author. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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metals Review

Metallurgical Challenges in Carbon Nanotube-Reinforced Metal Matrix Nanocomposites Abolfazl Azarniya 1 , Mir Saman Safavi 2 , Saeed Sovizi 1 , Amir Azarniya 1 , Biao Chen 3 , Hamid Reza Madaah Hosseini 1, * and Seeram Ramakrishna 4, * 1

2 3 4

*

Department of Materials Science and Engineering, Sharif University of Technology, Azadi Avenue, Tehran 11155-9466, Iran; [email protected] (Ab.A.); [email protected] (S.S.); [email protected] (Am.A.) Department of Materials Science and Engineering, University of Tabriz, Tabriz 51666-16471, Iran; [email protected] Joining and Welding Research Institute (JWRI), Osaka University, Osaka 567-0047, Japan; [email protected] Department of Mechanical Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore 117576, Singapore Correspondence: [email protected] (H.R.M.H.); [email protected] (S.R.); Tel.: +98-922-451-4107 (H.R.M.H.)

Received: 14 July 2017; Accepted: 11 September 2017; Published: 22 September 2017

Abstract: The inclusion of carbon nanotubes (CNTs) into metallic systems has been the main focus of recent literature. The aim behind this approach has been the development of a new property or improvement of an inferior one in CNT-dispersed metal matrix nanocomposites. Although it has opened up new possibilities for promising engineering applications, some practical challenges have restricted the full exploitation of CNTs’ unique characteristics. Non-uniform dispersion of CNTs in the metallic matrix, poor interfacial adhesion at the CNT/metal interface, the unfavorable chemical reaction of CNTs with the matrix, and low compactability are the most significant challenges, requiring more examination. The present paper provides a broad overview of the mentioned challenges, the way they occur, and their adverse influences on the physicomechanical properties of CNT-reinforced metal matrix nanocomposites. The suggested solutions to these issues are fully addressed. Keywords: carbon nanotube; nanocomposite; dispersion; interfacial adhesion; phase transformation; physicomechanical properties

1. Introduction Carbon nanotubes (CNTs) as one of carbon allotropes have been discovered in 1991 by Iijima [1]. There are two main types of CNTs including single-walled (SWNTs) and multi-walled nanotubes (MWNTs). CNTs are the monolithic cylinders of one or more layers of the graphene, so that if only one graphene layer is rolled, it is called SWNTs. MWNTs can be also understood by rolling several parallel graphene layers. Both SWNTs and MWNTs are formed with open or closed ends [2]. The morphology of MWNTs highly depends on their fabrication process, where Russian doll tubes are formed using perfectly concentric cylinders, and scroll tubes can be formed when a single graphene sheet is rolled as scroll [3]. All carbon atoms in the perfect structure of CNTs (except those present on the edges) are bonded in a hexagonal lattice. The presence of impurities drastically degrades the final properties of the tubes. MWNTs have larger diameters than SWNTs, where the diameter of SWNTs ranges from 0.8 to 2 nm and the diameter of MWNTs varies between 5 and 20 nm [2]. Furthermore, the size and surface area of CNTs are of prime significance, so that they can play an important role in the antibacterial applications, where a decrement in CNTs size may significantly improve their antibacterial efficiency

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Metals 2017, 7, 384

through the increase of the specific surface area [4]. Also, the variation of geometrical dimensions in CNTs can significantly affect their final properties due to a change in the length to diameter ratio. However, the final properties of SWNTs and MWNTs do not change in the same manner with the aspect ratio variation. For instance, the Young’s modulus of SWNTs deeply depends on the size and chirality, while the elastic modulus of MWNTs slightly changes with the diameter variation [5]. Although CNTs are among the strongest materials in nature, their mechanical properties vary in different directions. The results confirm the higher Young’s modulus and tensile strength in the axial direction compared to the radial one [6]. Owing to their appropriate mechanical, electrical, and thermal properties, CNTs have been extensively used in a wide range of the engineering applications such as high-strength and conductive composites, energy storage devices, and hydrogen storage [7]. Moreover, due to a favorable combination of strength and weight, CNTs are commonly used in the aerospace and automobile industries [8]. In these cases, CNTs are used whether in pure state or as a reinforcement in the metal, ceramic, and polymer matrix nanocomposites [9–11]. As to CNT-reinforced nanocomposites, the aim is the development of a new property or improvement of present properties. In general, the incorporation of CNTs into metallic matrices improves the mechanical properties of such binary systems (e.g., microhardness and fracture toughness) or electrical properties. However, they may adversely affect the properties [12]. Despite the mentioned advantages, this approach mainly faces four major challenges as follow: (i)

Favorable dispersion of CNTs throughout the matrix. A composite with improved properties will be obtained when the reinforcements are uniformly distributed through the matrix. Otherwise, the micro-pores as well as agglomerated particles may form all over the microstructure [13,14]. To overcome this challenge, a broad spectrum of dispersion methods is developed. Among these techniques, the mechanical methods [15], surface treatment [16], and chemical methods [17] are the most conventional. Each of these techniques has their own advantages and disadvantages, being discussed in the next sections. (ii) Unfavorable chemical reaction of CNTs with matrix at high pressures, elevated temperatures, and induced strains. It is usually accompanied by the formation of defects. On the other words, to thermally decompose CNTs in exposure to a metallic matrix, the presence or formation of defects is required. It is shown that the thermal decomposition of CNTs can bilaterally affect the final properties of the nanocomposites. In other words, the final properties strongly depend on the chemical composition of the formed intermetallics [18,19]. (iii) Poor interfacial adhesion between CNTs and the matrix due to the hydrophobic nature of the CNTs. This shortcoming deteriorates the load bearing between the matrix and CNTs. Moreover, intensive phonon scattering arisen from the insufficient adhesion can significantly enhance the electrical resistivity [20]. It is noteworthy that the interfacial adhesion can be improved whenever a controlled superficial chemical reaction between CNTs and metallic matrix occurs. On the other words, the poor interfacial adhesion can be considered as a sub-challenge under the title of “chemical reaction of CNTs with metallic matrix”. This review paper has adopted such a policy. (iv) Low compactability of metallic powders. The incorporation of CNTs into metallic powders can decline the relative density of final CNT-metal compacts, if the agglomeration of CNTs is heavy and their volume fraction is exceedingly high. As a general conclusion, the aforementioned challenges should be considered simultaneously in order to produce a CNT-reinforced metal matrix composite with superior properties. The present paper provides a broad overview of these challenges and their potential effects on the properties of CNT-dispersed metal matrix nanocomposites in some details. Also, the practical solutions to these challenges are introduced and their positive influences on the physicomechanical properties of the nanocomposites are studied. It is noteworthy that the full review of the potential effects of CNT

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addition on the physicomechanical properties of metallic systems is not the main purpose of the present review paper. However, one can find such information in [21,22]. 2. Metallurgical Challenges in CNT-Metal Matrix Nanocomposites 2.1. Dispersion of CNTs (Carbon Nanotubes) To date, a wide variety of research works are conducted with the aim to produce CNT-reinforced metal matrix nanocomposites with superior physicomechanical features for a broad range of functional and structural applications. However, the incorporation of CNTs into the metallic matrices is severely restricted due to both outstanding challenges: (i) Non-uniform dispersion of CNTs throughout the metallic matrices, and (ii) weak interfacial adhesion between CNTs and metallic matrix. In comparison to carbon particles and fibers, CNTs are more likely to be non-uniformly distributed inside the metallic systems due to their comparatively high aspect ratio and extraordinary specific surface area [23]. 2.1.1. Water Solubilization of CNTs A majority of electronic, thermal, and optical applications in which CNTs are used for the fabrication of novel devices or development of new emerging platforms need the large-scale production of stable CNT suspensions with a uniform dispersion and no agglomeration [24]. It is while the van der Waals attraction between CNTs results in their agglomeration and the formation of bundles (Figure 1) due to high surface area [25]. To overcome this practical challenge, some solutions are suggested. Better dispersion of CNTs in water as an important medium for biomedical and biochemistry applications is recognized by various methods such as using surfactants, polymers or chemical functionalization [24]. As to the chemical functionalization, a covalent bond is established between the chemical species and carbon atoms located at side walls or end caps of CNTs. However, this method suffers from two main problems [23,26–28]: (i)

(ii)

CNTs may shorten, get severe damages during the functionalization process, and lose their superior mechanical properties and distribution of π-electrons. As a general rule, the electron transport strongly depends on the surface structure of CNTs, because the structural damages can serve as phonon scattering sites and interrupt the transport properties of nanotubes. The strong acid solutions normally used for functionalization are hazardous to environment and hard to handle.

Figure 1. SEM (scanning electron microscope) images of (a) multi-walled carbon nanotube (MWCNT) bundles and (b) agglomerations [29] (Reproduced with permission from [29], Elsevier, 2012).

The aforementioned disadvantages have persuaded the researchers to develop other dispersing methods based on non-covalent modifications to uniformly distribute CNTs in water. These techniques benefit from two advantages over covalent functionalization methods: 5

Metals 2017, 7, 384

(i) (ii)

They induce no damages to CNTs, so that the electron transportation along the nanotubes remains intact; and CNTs are capable of forming an ordered network through the supermolecules acting as non-covalent modifiers. Among these modifiers, polymers and surfactants are the most conventional. These agents can thread themselves onto CNTs (Figure 2a) or wrap themselves around them (Figure 2b) [28,30,31]. Surfactants are also employed for dispersing CNTs in metal matrix nanocomposites.

Figure 2. Non-covalent functionalization of carbon nanotubes (CNTs): (a) With a surfactant and (b) with a polymeric agent [32] (Adapted with permission from [32], Elsevier, 2016).

2.1.2. Dispersion of CNTs into Metallic Systems The practical results show that the dispersion process of CNTs inside the metallic matrices can be completely different and more difficult than water media. This is because an intrinsically different phase (i.e., metal) is added to the system which makes the process more complex. Among these complexities, the CNT/metal interfaces and the poor wettability of CNTs by metal matrices are of prime significance. As generally accepted, an increment in CNT content can result in further agglomeration due to increased possibility of contacts between CNTs and decreased wettability. Therefore, dispersing individual CNTs in metallic systems with the reasonable spacing is a critical challenge for engineers and researchers [10,25]. The agglomeration of CNTs decreases the superior strengthening effects of nanotubes in CNT-metal matrix nanocomposites and deteriorates their physicomechanical properties. Moreover, the inhomogeneous dispersion of CNTs gives rise to non-homogeneity and anisotropy in these binary systems. On the other hand, it may impede the full sintering and reduces the relative density, wettability of CNTs, and resultant mechanical properties [16,33–35]. The above-mentioned challenges about the agglomeration and poor wettability of CNTs in metallic matrices have persuaded the researchers to discover more-efficient methods to uniformly distribute the nanotubes in these material systems. Their huge explorations have resulted in a wide 6

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variety of practical dispersion methods, although they are not classified into generally accepted standard categories. In fact, some literature has suggested different arbitrary classifications of the newly developed dispersion methods. For instance, they can be classified into four main groups: (i) Pre-treatments such as functionalizing, ultrasonication, and surface treatments; (ii) Mixing CNT with a metal precursor such as electrodeposition; (iii) Mixing CNT with a metal powders through milling prior to the consolidation routes; and (iv) Post-treatments such as extrusion and rolling. There are some different approaches for categorization of CNT dispersion methods. From the phenomenological perspective, they may be classified into three main groups: (i) Colloidal mixing in which physical reactions occur between CNTs and other species present in aqueous and inorganic media; (ii) Chemical mixing through which CNTs react chemically with other components; and (iii) Mechanical mixing which uses a mechanical force to detach CNTs from each other. However, all the categorized groups have no distinct boundaries, so that one dispersing method can belong to two or three groups in the same time. For instance, the ultrasonication method, which usually belongs to colloidal mixing, can be categorized in the chemical mixing group, if a chemical reaction occurs during the sonication of CNTs in the aqueous medium. As another example, the metallization is considered as a chemical mixing method, but it needs a complementary mixing method such as milling. As a result, the researchers have their arbitrary classification principles which are not general rules. Furthermore, CNT-metal dispersion methods are progressively updated and modified, so that the sharp boundaries between them are gradually faded. Table 1 summarizes some of the common dispersion methods among which the milling, ultrasonication, the application of surfactants, and metallization are explained in the next sections. Also, their possible challenges and suggested solutions are extensively discussed. Ball Milling In this procedure, a milling container with several hard balls is used for dispersing CNTs inside the metallic systems. During the milling, metal powders become fractured and welded again, so that CNTs are entrapped between them [15,36]. Since the primary powders as well as CNT bundles are entrapped between balls or balls and the container wall, the grinding motion of milling balls separates the nanotubes from each other, destroys the agglomerations, and disperses CNTs among the metallic powders. The milling time is one effective processing factor determining the degree of dispersion. As the mixing proceeds, the agglomeration is alleviated. Moreover, in the initial steps of milling, metallic powders have round shapes, and CNTs are dispersed at their surfaces. However, as the mixing continues, soft metallic powders flatters, and CNTs are embedded in the powders [15,36]. Generally, the mechanical milling is less able to inhibit CNT agglomerations, as seen in Figure 3 [37].

Figure 3. CNTs agglomeration in 5 wt % CNT-Al blends processed by mechanical milling for 30 min [38] (Adapted with permission from [38], Elsevier, 2011).

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Mass production Commercial applications Easy processing The possibility of in-situ functionalization or SDS (soduim dodecyl sulfate) treatment during the mixing Low cost Mass production Commercial applications Easy processing The possibility of in-situ functionalization or SDS treatment during the mixing Low cost

-

-

High efficiency (depending on the dispersion method used for metalized CNTs) The possibility of uniformly dispersing CNTs up to 5 vol % Mass production Production of CNT-containing systems with no damage The possibility of uniformly dispersing high CNTs content (up to 10–15 vol %) in metallic matrices The possibility of embedding CNTs inside the metallic powders rather than dispersing them at particles surfaces Mass production The possibility of automatization The possibility of uniformly dispersing a high CNTs content in metallic systems

-

-

-

-

In this method, CNTs are coated with a metallic layer such as Cu, Ni, Co, Mo, and W before their introduction into the matrix.

In this method, oppositely charged CNTs and metallic powder are co-deposited. CNTs and metallic powders are charged through a chemical step such as ultrasonication in acidic media.

A chemical or physical reaction occurs between functionalized CNTs and metallic ions in a solution medium.

In this method, CNTs are directly synthesized from the vapor phase on the metallic powders.

Surfactants

Metallization (i.e., decoration with metals)

Hetero-agglomeration principle

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Molecular level dispersion

In-situ chemical synthesis

-

-

-

-

In this method, non-covalent surfactants are embedded on CNTs surfaces with the aim to decrease the CNTs affinity to stick together.

-

In this method, CNTs are dispersed in a solution composed of organic solvents or aqueous surfactant by using high frequency sound waves.

In this method, CNTs and metallic powders are milled in a container with hard balls.

Ball milling

Advantages Mass production Commercial applications Easy processing Low cost Possibility of in-situ functionalization during the mixing The possibility of in-situ alloying

-

Ultrasonication

Principles

Dispersion Method

-

-

-

-

-

-

-

-

-

Need for expensive equipment

Practical complexity Limited production scale

Relatively high cost

Need for an additional mixing method such as mechanical mixing for producing dense parts

Low CNT content can be uniformly dispersed Need for a mechanical method to disperse CNTs throughout the metallic powders

Generation of defects in CNTs (lower than the mechanical milling) Low CNT content is homogeneously dispersed

Generation of defects in CNTs Cutting or breaking of CNTs Inclusion of contaminants into the system Long-time procedure

Disadvantages

Mg [56], FeCr [57], Cu [58], AlCu [59], Al [60], and Fe [61]

Cu [49,50], Ag [51], Ni [52], AgPd [53], Al [54], and AlCu [55]

Al [17]

Al [16] and Cu [48]

Al [45], Ti [18], Ni [46], and their alloys such as CuZr and CuZrAl [47]

Al [14], Cu [43], Fe [44] and their alloys

Al [39], Ni [40], Mg [12], Cu [41], Fe [42], and their alloys

Studied Metallic Systems

Table 1. A short description of CNTs (carbon nanotubes) dispersion methods in metal matrix nanocomposites, their advantages and disadvantages.

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9

In this method, CNTs are embedded in an electrochemically formed metallic film. This film is deposited as metallic ions transfer from cathode or electrolyte solution to the anode.

In this method, CNTs are added to other components such as SiC, A2 O3 , and graphite for better distribution. The ball milling is the most conventional method for mixing the hybrid reinforcements.

In this method, an initial preform of CNTs and other components is fabricated, and the molten metal is infiltrated into it.

After the fabrication of bulk CNT/metal nanocomposites, friction stir processing (FSP) or selective laser melting (SLM) is used for re-dispersing CNTs.

In this method, a solid metal with a CNT-containing flux is poured into an induction furnace. The presence of such a flux decreases the melting temperature of the solid metal. Finally, CNTs are dispersed by induction-formed fluctuations.

Electroless deposition

Electrodeposition

Using hybrid reinforcements

Preform infiltration

FSP/SLM induced dispersion

Induction melting-based dispersion

In these methods, CNTs are dispersed in a molten metal through stirring or squeezing.

In this method, CNTs are immersed in an electrolyte solution containing metal ions. During the process, the nanotubes are covered by ions through developing potential. A substrate can be used on which CNTs and metallic ions are simultaneously deposited.

Casting-based methods such as stir casting and squeeze casting

Principles

Dispersion Method

The possibility of obtaining higher mechanical properties due to strengthening effects of hybrid reinforcements High efficiency Cost-effective The possibility of selectively dispersing CNTs at a macro level High efficiency to prevent the agglomeration

-

-

Mass production Easy to handle

Mass production Easy to handle Comparatively low cost

-

-

-

-

-

-

High efficiency (depending on the dispersion method used before the fabrication)

Commercial applications High efficiency The possibility of combination with other mechanical mixing techniques such as ultrasonication for better CNTs dispersion The possibility of uniformly dispersing other components

-

-

Commercial applications High efficiency The possibility of combination with other mechanical mixing techniques such as ultrasonication for better CNTs dispersion The possibility of uniformly dispersing other components

-

Advantages

Table 1. Cont.

-

-

-

-

-

Additional cost due to the melting of the metals Comparatively low efficiency

Additional cost due to melting of the used metal Need for expensive instruments

Imposition of additional cost Only used for finished CNT/metal bulk nanocomposites

Metallic component needs to be melted Restrictions imposed by shape geometry Need for a two-step procedure Relatively high cost

-

-

Need for an additional mixing process such as ball milling The possibility of formation of unwanted intermetallic compounds

Only used for certain metals High cost Not easy to handle Need for a supporting mixing method to obtain a full efficiency

Only used for some metals High cost Need for a supporting mixing method to obtain a full efficiency

-

-

-

Disadvantages

Al [80]

Al [78], Ti [79]

Al [76], Cu [77]

Al [74,75]

Al [71], Cu [72], Mg [73]

Ni [66], Cu [67], Sn-Bi [68], Cu-Ni [69], and Ni-Co [70]

Cu [62], Ni [63], Al [64], Au, and Pt [65]

Studied Metallic Systems

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To improve the CNT dispersion in metallic matrices, a process controlling agent (PCA) such as polyacrylic acid (PAA) or methanol is usually used. These chemicals are utilized not only for better dispersion, but also for CNTs functionalization (Figure 4).

Figure 4. TEM (transmission electron microscopy) images of (a) as-synthesized CNTs, (b) CNTs ball milled with NH4 HCO3 as a functionalizing agent and process control agent (PCA), and (c) CNTs milled with no PCA. The presence of chemicals is shown to result in improved CNTs dispersion [81] (Reproduced with permission from [81], Elsevier, 2008).

Despite the ability to functionalize CNTs, the mechanical milling is capable of in-situ alloying as a subsidiary mechanism for strengthening CNT-metal matrix nanocomposites. Moreover, it can produce the metallic nanostructures whose mechanical strength is higher than that of conventional nanocomposites. Such grain structures arise from the formation of large numbers of dislocations inside the metallic powders due to high plastic deformation. As another advantage, a close control over the particle size and particle size distribution in ball milled powder systems is possible. These characteristics are controlled by two competing mechanisms during the milling: (i) Cold working, which results in decreased ductility, particle fracturing, and reduced particle size; and (ii) Cold welding, which results in increased particle size [40,82,83]. Besides the aforementioned advantages, the mechanical milling may induce structural damages to CNTs such as breaking or shortening, especially when the dwell time is extended. These defects confine the processing time [84]. They may be also formed when the milling is conducted under inappropriate practical conditions. The non-controlled milling results in formation of carbides through a chemical reaction between the damaged nanotubes and reactive metallic powders such as Al. Such reactions may occur even at comparatively low milling temperatures. Among the various milling apparatuses, the planetary milling imposes less energy on metallic powders and CNTs, especially in comparison with high-energy mechanical milling. The higher the energy imposed on the powder blend, the higher the volume fraction of defective nanotubes will be [82]. In contrast, high energy milling can evenly embed CNTs in the metallic powders, while the planetary milling is usually incapable of uniformly dispersing CNTs. It may persuade the CNT agglomerations and locate them at the external surface of metallic powders rather the diffusion into them. Furthermore, a useful combination of practical conditions is required to obtain high-quality nanocomposites [85]. Figure 5 shows the way the nanotubes may be mechanically dispersed throughout Al matrix. It seems that the utilization of a chemical modification or a processing control agent as well as the low energy milling technique can successfully reduce the structural damages of CNTs [82]. Another practical method to prevent the chemical reactions during the milling is the suppression of heat-induced changes in metallic powders through cryogenic techniques. Cryogenic milling is referred to as the milling procedure at very low temperatures (usually below 100 K). Such temperatures give rise to the increase in brittleness of metallic powder, reduction in grain growth, resultant grain refinement, and decrement in grain size. In this case, the hardness of CNT-metal matrix nanocomposites increases by the Hall-Petch mechanism. Therefore, such methods are able to reduce the processing 10

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time and bring forth finer grains with the lowest possibility of chemical reactions. Interestingly, if the cryogenic milling is conducted with the presence of carbon-containing PCA such as stearic acid, their carbons may react with metallic powders and form carbides during the milling or subsequent sintering processes [82].

Figure 5. SEM images of a CNT-Al blend after milling by high energy milling (a–d) and low energy milling (e,f) [85]. The CNTs are uniformly dispersed and embedded in Al powders during high energy ball milling, as shown by yellow arrows (Reproduced with permission from [85], Elsevier, 2011).

The CNTs are uniformly dispersed and embedded in Al powders during high energy ball milling. However, their dispersion has not a uniform pattern, so that there are some Al powders with no CNT. Conversely, all Al powders may obtain CNTs during the milling by low energy milling, but CNT agglomerations still remain, and CNTs are located at the surface of Al powders [85]. In summary, if the best equipment as well as better control over processing conditions is utilized to disperse CNTs among the metallic powders, each method has its own intrinsically certain ability, so that if the CNT loading exceeds the optimal value of each technique (4–5 wt % CNT at the best conditions), the nanotubes tend to form agglomeration. In this case, the pores and poor interfacial bonding can degrade the mechanical properties as well as thermal and electrical characteristics of manufactured CNT-containing metal matrix nanocomposites [86,87]. Consequently, many researchers believe that such a dispersion method and arbitrary practical conditions have low potential to uniformly disperse CNTs in metallic matrix. Some pros and cons of mechanical milling are listed in Table 1. Ultrasonication In this method, sound waves with high frequency are used for dispersing CNTs in a solution composed of organic solvents or aqueous surfactants. In order to increase the solubilization of CNTs in

11

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water, the hydrophobic CNT surfaces can be modified by a wide range of surfactants and polymer adsorbates. The mechanism by which the ultrasonication debunds CNTs in a solution containing an aqueous surfactant is composed of four step: (i) A high local shear stress at the end of CNT bundles is induced by the collapse of micro-bubbles formed during the ultrasonication; (ii) Surfactants diffuse into and absorb on CNTs as soon as a gas bubble or gap is formed; (iii) The unzipping process takes place along the longitudinal axis of a particular CNT; (iv) The individual surfactant-coated CNT is released inside the solution [88]. It is worth mentioning that unzipping the individual CNTs from the bundles will not be completed if there is no polymer absorbate or surfactant. This is because CNTs have van der waals attraction to each other which may again [88]. However, the formation of hotspots during the ultrasonication which simultaneously increases the local pressure and temperature may induce some defects in CNTs, form carbon-dangling bonds at their sidewalls, and cut or chop the nanotubes. Such drawbacks decrease the effectiveness of the ultrasonication as a useful method for dispersing CNTs in CNT-containing metal matrix nanocomposites. In order to reduce these disadvantages, organic molecules, for example, poly (methyl methacrylate) (PMMA) or monochlorobenzene (MCB) can be added to form reactive species during the sonication. These species which are generated from decomposition of added organic molecules at the hotspots can chemically react with dangling bonds of CNTs and are easily wiped through burning in oxygen gas. This event may leave holes on the sidewalls [89]. As a holistic conclusion, the ultrasonication has a finite ability to disperse CNTs in metal matrix nanocomposites. According to the literature, the acid-treated CNTs can be uniformly dispersed up to 1 wt % in metallic matrices (such as Al) by this technique. Further addition of CNTs to the matrix results in agglomeration and generation of micropores during the subsequent consolidation stage. These defects may significantly decrease the mechanical, electrical and tribological properties of these binary systems [13,14]. The Application of Surfactants Active surface agents such as surfactants can significantly improve the surface conditions of CNTs and prevent from their agglomeration and re-bonding. In general, these surface agents consist of two parts: (i) A hydrophilic part with a polar head and (ii) A hydrophobic part coupled with a hydrocarbon chain at the tail of the surfactant molecule. These modifiers enhance the metastability of the colloidal suspensions by providing an electrostatic and/or steric repulsive force between CNTs and reducing their surface energy. Depending on their head group charges, the surfactants are usually classified into three different groups: cationic, anionic, and nonionic or zwitterionic [90]. One of the most conventional anionic surfactants for CNT-reinforced metal matrix nanocomposites is SDS (soduim dodecyl sulfate), wherein the sulfates and hydrocarbon chains act as hydrophilic and hydrophobic parts, respectively. In fact, CNTs can interact with the hydrophobic part of SDS, so that an electrostatic repulsion force is formed between the negatively charged CNTs and the hydrophilic sulfate groups of SDS. Thus, it can remarkably affect the stabilization of CNTs in aqueous media [28]. There are quite a few studies evaluating the application of SDS as a surfactant in metallic matrices. The surveys confirm that there is a defined limit for inclusion of CNTs into the matrices. This is because if the critical level is exceeded, the uniform dispersion of CNTs throughout the metallic matrices is not possible and further incorporation leads to the formation of agglomerated particles [45]. The agglomerated particles can impressively degrade the mechanical properties and functionality of the nanocomposites. However, the introduction of SDS into the metallic systems can improve the CNTs dispersion. Zwitterionic surfactants are a new class of widely used surface agents for dispersing CNTs in metal matrices. Whereas they are polar and soluble in water, their solubility in organic media is limited. They chiefly consist of cationic and anionic parts attached to the same molecule. The self-assembled monolayers (SAMs) of these surfactants can be also formed on CNT bundles (Figure 6). In this case,

12

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the strong dipole/dipole attraction between SAMs and CNTs overcomes the present van der Waals forces between the individual nanotubes and disassembles the ubiquitous bundles [91,92].

Figure 6. A schematic illustration of the mechanism by which zwitterionic surfactants disassemble CNT bundles [47] (Reproduced with permission from [47], Elsevier, 2014).

The experimentations show that if the CNTs content exceeds 1 wt %, the agglomerated particles will form all over the microstructure of the nanocomposites. Moreover, the processing parameters deeply affect the way the nanotubes are dispersed throughout the matrix, where the mechanical properties can be degraded because of employing an unsuitable process. It is ascribed to the formed pores and degraded interfacial adhesion due to evaporation of the surfactant at high temperatures. Post treatments (e.g., heat treatment) is often employed to remove such pores [92]. Another disadvantage is the lack of control on the reactions between CNTs and the metal matrices, where the intermetallics can be generated as a result of unwanted reactions during the sintering process [92]. It seems that there is an urgent need for other surface treatments such as coating to uniformly disperse the high volume fraction of CNTs throughout the metal matrices. Metallization The outermost surface of CNTs can be coated by several metallic elements such as Cu, Ni, Co, Mo, and W, if enhanced interfacial adhesion between CNTs and the matrix is greatly required. Such coatings may contribute to the prevention of CNTs agglomeration all over the matrix [93,94]. It is demonstrated that W-coated CNTs by metal-organic chemical vapor deposition (MOCVD) method can significantly inhabit the CNTs agglomeration. This fact is shown in Figure 7. As clearly seen, the pores or air bags at CNT/metal interface is removed as the main reason for improved interfacial adhesion [48]. The enhanced thermal conductivity of the composites is ascribed to the formation of desirable interface which can remarkably reduce the thermal lost. Similarly, it is shown that there is no intermetallic formation in the microstructure of (Mo-coated CNT)-Al nanocomposites [16]. It is attributed to the presence of Mo layer, acting as an appropriate barrier to the interfacial reaction between CNTs and the matrix. Additionally, the applied coating is able to impressively increase the volume fraction of the uniformly dispersed CNTs. The mechanical properties of these composites are proved to improve due to desirable dispersion of CNTs throughout the matrix. Despite the improved mechanical behavior, the exploitation of such coatings may slow down the electron movement all over the nanocomposites and deteriorate the electrical conductivity.

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Figure 7. SEM micrographs of (W-coated CNT)-Al nanocomposites fabricated by ultrasonication method: (a) 5 vol % CNTs, (b) 10 vol % CNTs, (c) TEM micrograph of (W-coated 5 vol % CNT)-Al nanocomposites, and (d) higher magnification of the selected area in (c) [48] (Reproduced with permission from [48], Springer, 2012).

Effects of Dispersion Methods on Physicomechanical Properties In addition to the consolidation and sintering techniques, the dispersion methods may greatly influence the properties of CNT-containing metal matrix nanocomposites. A review on the literature confirms that whenever the used dispersion method is inappropriately selected, so that it keeps or forms CNT aggregates, the mechanical and functional properties of produced nanocomposites may be degraded due to the formation of pores and poor interfacial boundaries between nanotubes and metallic matrix. Figure 8 shows the hardness of some spark plasma sintered (SPSed) CNT-Al nanocomposites as a function of CNT content for different dispersion methods. The aim behind the selection of SPS as the sintering method and Al as the metal matrix is to eliminate the potential effects of sintering process and chemical composition of metal matrix on the properties of these systems as much as possible. As clearly seen, the highest hardness values can be achieved by the mechanical milling due to the formation of severe plastic deformation-induced dislocations and defects in Al matrix. However, if the dispersing process is not suitably controlled, its intrinsic potential will be obtained. For instance, some dispersion techniques result in decreased properties when the CNT content exceeds 1–1.5 wt %. As another example, if the metallization is not done in a suitable manner, or the coating layer is not truly selected for the metal matrix, it cannot enhance the mechanical properties. Such a justification is also applicable to ultrasonication. If the ultrasonication is carried out under non-controlled conditions, the CNTs cannot play their positive role and reduce the mechanical properties after adding a low CNT content. The tensile strength of MMNCs-reinforced metal matrix nanocomposites has similar trends. The incorporation of CNTs increases the tensile strength of metallic matrices, if they are homogeneously dispersed and make a proper interfacial bonding with the matrix. However, by adding CNTs more than the critical content, a heavy agglomeration occurs and the tensile strength is degraded (Figure 9). Similar trends are also seen in the thermal conductivity of SPSed CNT-Cu or Al matrix nanocomposites (Figure 10). Generally, the addition of CNTs to the metallic matrices decreases the thermal conductivity due to the grain boundary-induced electron scattering. However, the metallization can increase such a property through enhanced interfacial boundaries between nanotubes and metallic matrix. Besides the metallization, the polymer wrapping and acid treatment can decrease 14

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the degradation of thermal conductivity of Cu matrix by adding CNT due to the formation of enhanced interfacial boundaries. Such methods also have a limited potential to disperse CNTs. If the CNT content exceeds an optimum limit, the agglomeration of CNTs will decrease the thermal conductivity of these binary systems.

Figure 8. Hardness of some spark plasma sintered (SPSed) CNT-Al matrix nanocomposites fabricated by various dispersion methods [13,16,36,95].

Figure 9. The tensile strength of some SPSed CNT-metal matrix nanocomposites fabricated by various dispersion methods [14,96,97].

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Figure 10. Thermal conductivity of some SPSed CNT-Cu and CNT-Al matrix nanocomposites fabricated by various dispersion methods [48,98–103].

2.2. Consolidation Challenge The consolidation of pure metallic powders and their alloys is not a critical challenge, and pure metallic parts with the full density can be obtained by various sintering methods such as hot pressing, hot isostatic pressing, and spark plasma sintering. However, the addition of CNTs to the metallic matrix may greatly affect the consolidation rate and relative density of the binary system. In fact, the poor wettability between CNTs and metallic matrices and the formation of CNT agglomeration-induced pores impede the densification and decrease the relative density of the nanocomposites; no matter what the sintering method is. As a result, producing a fully dense CNT-reinforced metal matrix nanocomposite is a critical challenge which highly depends on the dispersion of CNTs in the metallic matrix. Actually, the selection of a proper dispersion method capable of homogeneously dispersing CNTs can approach the relative density of the CNT-metal bulk systems to that of pure metals. Figure 11 shows the relative density of different CNT-reinforced metal matrix nanocomposites as a function of CNT content and dispersion method. As seen, the CNT ddition usually decreases the relative density and sinterabillity of the metallic matrices. Moreover, an increment in CNT content generally brings about more CNTs agglomeration and pores, decreasing the relative density. However, if a proper dispersion method such as the metallization for Cu and electroless deposition for Ag is utilized, CNTs will be homogeneously dispersed in metallic matrices with no declined density. It is noteworthy that each dispersion method has a certain ability to homogeneously disperse CNTs in a specific metallic matrix. If the CNT content exceeds the maximum ability of a mixing method, the agglomeration will occur. For instance, as seen in Figure 11, CNTs up to 2.5 vol % are homogeneously dispersed in Ag matrix through the electrodeposition method. Further an increase in CNT content may degrade the consolidation by inducing the agglomerates and pores. As a conclusion, the consolidation of CNT-containing nanocomposites highly depends on the uniformity degree of dispersed CNTs. It implies the importance of the dispersion method to achieve a fully dense part. The employment of a suitable dispersion method and addition of CNTs below the critical content may assure the manufacturing of fully dense CNT-reinforced metal matrix nanocomposites.

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Figure 11. The relative density of some SPSed CNT-metal matrix nanocomposites fabricated by different dispersion methods [16,48,49,96,104].

2.3. Chemical Reaction with Metallic Matrix 2.3.1. Thermal Decomposition of Pure CNTs A review on the literature confirms the effective function of CNTs on the structural and functional characteristics of CNT-containing bulk materials. To exploit the full potential of CNTs in nanocomposites, one should deeply understand the structural phenomena and microstructural interactions likely occurring during the synthesis or post-treatments of these new emerging materials. The first key to this issue is the profound consideration of structural transitions and phase transformations of CNTs in pure state and in exposure to a metallic matrix, because they can adversely affect the intrinsically superior properties of pure CNTs [105–109]. Table 2 summarizes the recently reported structural transitions of pure CNTs in a variety of densification methods and their influences on the properties. As seen, less empirical results are reported in the literature addressing these structural changes. In general, any allotropic transition in CNTs is highly undesirable, because it can limit the intrinsically unique features of nanotubes in electronic aspects and load bearing performances [110–112]. An overwhelming majority of reports about the conversion of CNTs to other allotropes is associated with spark plasma sintering, where high pressure and elevated sintering temperatures However, other persuade any thermally activated phase transformation [105,106,113,114]. consolidation techniques such as hydrogen plasma [115], high-pressure/high-temperature annealing heat treatment [116–118] and chemical vapor deposition (CVD) [107,108,119] can thermally decompose the structural geometry of nanotubes. As a practical conclusion, the engineers often strive to prevent from the thermal degradation of CNTs in both pure state and composite form, since it can vanish the final properties [120]. Hence, one can determine two temperature regimes for CNT allotropic phase transformations depending on whether CNTs are preserved: (i) low-power regime and (ii) high-power regime. In the first case, known as safe regime, the processing pressure and temperature are sufficiently low, so that the thermal power cannot destroy the integrity and tubular morphology of nanotubes and thermally activate the mechanisms by which CNTs can structurally convert to other allotropes. In the second case, the elevated temperature or high applied pressure can persuade the thermal decomposition of nanotubes. There is not any sharp temperature boundary between these regimes, depending on the kind of consolidation procedures, type of matrix exposed to CNTs, and activated sintering mechanisms [121].

17

1600

1000, 1500, 2000

Spark plasma sintering + polishing

Spark plasma sintering

18

1500

1700, 1800, 1900, 2000

1200

700

Spark plasma sintering

Spark plasma sintering

Spark plasma sintering + Fe35Ni catalysts

Chemical vapor deposition

Below 2000 2100–2400 2500–2800

1500

Spark plasma sintering

(i) (ii) (iii)

1500

Spark plasma sintering

Annealing heat treatment

Processing Temperature (◦ C)

Fabrication Process

-

70

100

80

-

50

60

80

80

Applied Pressure (MPa)

30

20

5

20

30

5

1

20

20

Processing Duration (min)

graphite

diamond + graphite

-

-

The presence of graphene in all targeted temperatures

Partial conversion of CNTs into diamond The formation of n-diamond

-

-

No structural changes the formation of large-sized double-walled carbon nanotubes (DWCNTs) the graphitization of multi-walled nanotubes (MWNTs) and other kinds of crystalline sp2 hybridized carbon structures

-

graphene

graphene

diamond + graphite

diamond

Final Phase Structure

Not reported

The formation of diamond through a layer-by-layer mechanism

The formation of graphene nanosheets due to the evaporation of carbon atoms during the heating stage of SPS and their agglomeration during the cooling stage

High localized temperatures arisen from the presence of plasma during the sintering as a prerequisite for the formation of diamond

Not reported

Not reported

Not reported

An increment in electrical conductivity due to the formation of graphene nanosheets

Not reported

[107]

[124]

[123]

[122]

[116]

[114]

Enhanced thermal stability due to the formation of graphene at a high sintering temperature

Not reported

[106]

A significant reduction in friction coefficient due to the much lower friction between formed graphene layers

Shear stress induced by polishing which leads to peeling graphene away from CNTs

Not reported

[113]

[105]

Ref.

Not reported

Not reported

Reported Changes in Properties Due to CNTs Structural Changes

Not reported

Not reported

Reported Mechanisms for Structural Changes

Table 2. Overview over a variety of conventional manufacturing techniques, allowable processing windows and structural changes of pure CNTs and their effects on the physicomechanical properties of nanotubes.

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19

726

Hydrogen plasma

1600–2000

-

Electrical breakdown method

High pressure and high temperature heat treatment

500

Chemical vapor deposition

1100–2000

1300

Annealing at high pressure

High temperature annealing heat treatment

2226

Compaction + laser beam heating

55–70

700

Chemical vapor deposition

Solution-based oxidative process

Processing Temperature (◦ C)

Fabrication Process

≥15,000

-

-

-

-

-

4500

17,000

-

Applied Pressure (MPa)

graphite nanoribbons

diamond

10–10,000 s (Exact dwelling time has not been reported)

graphene nanoribbons + graphite oxide

diamond

graphene nanoribbon

highly ordered graphite

diamond

diamond

amorphous carbon

Final Phase Structure

240 min heating + 480 min cooling process

60

-

-

2160

12

-

56

Processing Duration (min)

Table 2. Cont.

Not reported

Not reported

Fast unwrapping of MWNTs with no an intermediate step The formation and growth of diamond particles due to the formation of amorphous carbon clusters with sp3 bonds

The formation of diamond phase as a result of atomic diffusion with no graphitization or formation of intermediate phases

Not reported

Not reported

Not reported

CNT → carbon nanofiber → highly ordered graphite

The thermal activation of three different mechanisms for different heat treatment temperatures: The coalescence of SWNTs at 1400 ◦ C followed by rearrangement of the present bonds at 1600 ◦ C to form MWNT The collapse of MWNTs into graphite nanoribbons at temperatures higher than 1800 ◦ C Complete conversion of MWNTs to graphite nanoribbons at temperatures higher than 2000 ◦C

Not reported

The conversion of CNTs to quasi-spherical onion-like structures followed by the formation of diamond crystals

Not reported

Not reported

Direct conversion of CNTs to nano-sized diamonds with no an intermediate step such as melting or dissolution

The formation of graphene nanoribbons due to complete unravelling of CNT side walls

Not reported

Reported Changes in Properties Due to CNTs Structural Changes

Not reported

Reported Mechanisms for Structural Changes

[127]

[118]

[109]

[115]

[126]

[119]

[117]

[125]

[108]

Ref.

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In summary, CNTs may transform into other thermodynamically stable allotropes during the heat treatment process. Among these allotropes, cubic diamond [105], n-diamond [122], mono/multi-layer graphene [106], and graphite sheets [107] are the most extensively reported. However, the way these allotropes can affect the final properties are highly contradictory. Whereas some literatures refer to the deterioration of physicomechanical properties of CNT-based bulk materials, others put emphasis on the drastic improvement of some properties [106,123]. On the other hand, there is no consensus among researchers regarding what temperature range guarantees the conversion of CNTs to other carbon allotropes and which forms of new phases can be formed. For instance, Zhang et al. [123] evaluated the microstructure, mechanical and electrical properties of spark plasma sintered CNTs in pure state and demonstrated the in-situ formation of graphene nanosheets at 1700–2000 ◦ C. It was shown that the formed graphene nanosheets are responsible for improved electrical conductivity in these conditions. Figure 12 indicates the TEM image of this formed graphene nanosheet at 2000 ◦ C.

Figure 12. TEM image of graphene nanosheet formed during spark plasma sintering at 2000 ◦ C [123] (Adapted with permission from [123], Elsevier, 2013).

2.3.2. Chemical Reactions The chemical reactions between CNTs and metallic matrices can be categorized into two different aspects: (i) The partial chemical reaction at CNT/metal interface and (ii) the complete chemical reaction of CNTs with metallic matrix. If controlled, the former is favorable for the load bearing applications, but the latter can bilaterally affect the physicomechanical properties of CNT-dispersed metal matrix nanocomposites. In the next sections, the practical and scientific aspects of these chemical reactions will be addressed in some details. Complete Chemical Reactions (a) CNT-Reinforced Al Matrix Nanocomposites A review on the literature confirms that Al4 C3 is the most common in situ formed intermetallic in CNT-reinforced pure Al matrix nanocomposites. However, other intermetallics such as SiC can be formed during the reaction between the alloy matrix and CNTs. In fact, the chemical composition of the formed intermetallics strongly depends on the fabrication temperature and the chemical composition of the alloy matrix. The kinetics of the carbide formation at temperatures lower than Al melting point is negligible. Moreover, the presence of some alloying elements can significantly control the final properties of the nanocomposites. Two different sources are suggested for the formation of Al4 C3 : (i) Used process control agents (PCA) during the milling process such as stearic acid and (ii) incorporated CNTs. Also, the processing parameters such as temperature, post treatment variables 20

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and the size of initial CNTs play an important role in the formation of carbides and controlling their volume fraction, morphology, and geometrical dimensions. It is found that the morphology of Al4 C3 can be easily changed by applying the post heat treatment or the presence of initial defects in CNT structures. Four types of morphologies including whisker-like, dumbbell-like, needle-like, and tube-like ones can be observed for Al4 C3 particles, depending on the post treatment conditions and the purity of initial CNTs. Similarly, the volume fraction of the formed carbides deeply depends on the geometrical size of initial CNTs. Generally speaking, the presence of carbides at CNT/Al interfaces improves the mechanical properties of the nanocomposites. It is attributed to the enhanced load transfer between the matrix and reinforcements as a result of carbide formation at the interface. The strengthening effect of the formed carbides depends on their thickness, so that the formed carbides will dissolve back into the melt if the critical thickness is less than a critical limit. The in situ formation of Al4 C3 at the Al/CNT interface has been subjected to the extensive research studies. Some research works have considered the precipitation of Al4 C3 as one of the predominant strengthening mechanisms in CNT-dispersed Al matrix nanocomposites. This compound forms during a chemical reaction between Al matrix and CNTs as: 4Al + 3C = Al4 C3 ΔG = −289,512 + 60T, T < 660 ◦ C

(1)

The free energy (ΔG) of this reaction is negative by the melting point (660 ◦ C) of Al. It means that Al4 C3 is thermodynamically stable at temperatures lower than 660 ◦ C. Coincidently, the formation of this carbide at temperatures lower than 660 ◦ C is also reported [128]. The mentioned reaction can be explored for ball milled CNT-Al powder blend. In this case, Al4 C3 can be formed when the energy state of the mixed powders reaches to a sufficiently high level. It is believed that the long-term milling process can seriously damage the CNTs structure and lead to the in situ formation of Al4 C3 [87,129]. For instance, Ostovan et al. [87] evaluated the microstructure of the milled powders for 8 and 12 h. It was found that the needle-like Al4 C3 is just formed after the milling for 12 h. It was attributed to the long-term milling-induced damage in CNTs. As a basic conclusion, the possible CNT amorphization during the sintering process as well as the formation of crystallographic defects in CNTs can significantly facilitate the formation of Al4 C3 . It is ascribed to the synergetic effects of instability of milled powders after the long-term milling, and generated driving force as a result of thermal processing such as sintering [87]. The literature have suggested two different sources for Al4 C3 formation: (i) Used process control agents (PCA) during the milling process, (e.g., stearic acid) and (ii) incorporated CNTs. Since the stearic acid is composed of oxygen, carbon, and hydrogen atoms, it can be decomposed into its constituent elements during the long-term milling. On the other hand, the milling process can significantly damage the CNTs structure and enhance their amporphization [130]. The produced atomic carbon from the mentioned sources can easily react with the matrix Al to form Al4 C3 . More importantly, the formation of Al4 C3 may consume the CNT walls or the entire CNT [129]. If the alloy matrix is used, other intermetallics can be formed beside Al4 C3 . The chemical composition of the formed intermetallics strongly depends on the alloying elements. For instance, Al4 C3 and SiC can be formed as a result of the reaction between Al-Si alloy matrix and CNTs. Bakshi and coworkers [131] showed that the formation of each intermetallic corresponds to the Si weight percent at the fabrication temperature, as shown in Figure 13. As seen, the more the Si content in the alloy, the more SiC particles will form. It is ascribed to increased amount of Si in the liquid melt due to the reaction between Al and C to form Al4 C3 . It can increase the activity of Si to form SiC. Both reactions will proceed in the same manner until all of carbon content is completely exhausted.

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Figure 13. The chemical composition dependence of formed intermetallics on Si weight percent and fabrication temperature [131] (Reproduced with permission from [131], Elsevier, 2009).

In CNT-(Al-Si) nanocomposites, the poor wettability arises from a major difference between the surface tension values of the alloy matrix (about 800 mN·m−1 ) and CNTs (about 45 mN·m−1 ) [33]. However, the literature have reported improved wettability with the generation of some intermetallics. It is attributed to reduce contact angle between the alloy matrix and CNTs through the infiltration of Al-Si melt into CNT clusters [131]. The formation kinetics of Al4 C3 at temperatures lower than Al melting point is slow. For the binary composites in which CNTs react with Al matrix chemically, the presence of alloying elements can significantly control the final properties, because the chemical composition may affect the thermodynamic aspects of the Al4 C3 -forming reaction [132]. The processing parameters such as temperature, post treatment variables and the size of initial CNTs play an important role in the formation, volume fraction, morphology, and size of the formed carbides in CNT-Al matrix nanocomposites. These compounds can be precipitated at different locations. For instance, the nano-sized Al4 C3 particles often form on the external surface of CNTs as a result of a reaction between Al and CNT-derived amorphous carbon. In the case of partially graphitized CNTs, the carbides are evolved at the end of damaged CNTs. It is shown that only a little amount of MWNTs can react with the matrix to generate carbides and contribute to enhanced properties of the composites [133]. There are two main factors affecting the morphology of the formed compounds during the reaction between the matrix and CNTs: (i) Post treatment, and (ii) internal defects of CNTs. The morphology of Al4 C3 precipitates strongly depends on post heat treatment temperature. It has been reported that the carbides grow in the form of nanosized whiskers at 950 ◦ C, while they are more likely to disappear at comparatively lower temperatures [133]. Additionally, one can suppress the formation of these carbides through the close control over the sintering conditions, i.e., the sintering temperature should be kept below Al melting point [134]. Figure 14 exhibits the elongated Al4 C3 precipitates formed as a result of a high-temperature heat treatment. As obviously seen, the longer the heat treatment temperature, the more elongated the carbides are. In addition to the processing parameters, the internal defects of nanotubes can also change the way the precipitates are evolved. Two fascinating types of morphologies are reported for Al4 C3 particles, i.e., dumbbell-like and tube-like. Whereas the first morphology arises from the CNTs tips, the second case originates from defective CNTs [135]. Figure 15 indicates SEM images of grown nano-whiskers during annealing treatment at 950 ◦ C.

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Figure 14. HRTEM (High resolution transmission electron microscope) images of MWNT (multi-walled nanotubes)-reinforced Al matrix nanocomposites heat treated at: (a) 873 K for 6 min and (b) 883 K for 60 min [132] (Adapted with permission from [132], Elsevier, 2016).

Figure 15. SEM images of (a) grown nano-whiskers during annealing treatment at 950 ◦ C, and (b) the formation of nano-whiskers on the Al grains, as shown by the white arrows [133] (Adapted with permission from [133], Elsevier, 2006).

The geometrical size of initial CNTs is one of the affecting parameters controlling the volume fraction of the formed carbides. The empirical results show that Al4 C3 particles can easily arise from MWNTs with 40 nm in diameter, while the amount of Al4 C3 seems to decrease with an increase in diameter by 140 nm examined from the XRD results [38]. Albeit the development of Al4 C3 carbides in CNT-Al matrix nanocomposites is of prime significance, the way they can affect the microstructure-related and mechanical properties is more important. As generally believed, the presence of carbides at CNT/Al interfaces improves the mechanical behavior of these nanocomposites. This is because the carbides can enhance the matrix capability to efficiently transfer the external load to nanotubes. This approach is in good agreement with the results reported by Tjong et al. [136], Esawi et al. [137], Kwon et al. [130,135] and Chen et al. [138]. Table 3 gives some information about the presence of Al4 C3 in some CNT/Al composites with high strengthening effect. The results indicate that the formation of Al4 C3 maybe doesn’t have an adverse effect on the composite strength. The recent study by Chen et al. [138] reveals that nano-sized interfacial Al4 C3 can noticeably enhance the load transfer efficiency, resulting in improved composite strength. The way the intermetallics affect the final properties of the nanocomposites strongly depends on their thickness. The critical thickness is estimated as [131]: tcri = −V M × (Δγ/ΔGf )

23

(2)

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where V M is the molar volume of the formed carbide, Δγ is the increase in total surface energy due to newly formed interfaces, and ΔGf is the free energy of formation per mole of carbide. Further growth of the formed carbides can be achieved when the thickness reaches to a critical value. It can remarkably improve the wettability through diminishing the contact angle [131]. Table 3. The variation of tensile strength in CNT-dispersed Al matrix nanocomposites with the volume fraction of Al4 C3 precipitates (SPS: Spark plasma sintering). Amount of Added CNTs

Fabrication Method

Increase in Tensile Strength Value (%)

Ref.

2 wt % MWNTs 5 vol % MWNTs 5 wt % MWNTs

Sintering → hot extrusion SPS → hot extrusion Sintering → pressing → annealing

57.5 128 184

[139] [120] [128]

The in situ carbides evolved in CNT-(Al-Si) nanocomposites grow through two different pathways: (i) At the Al4 C3 /Al-Si interface and (ii) at the Al4 C3 /CNT interface. In accordance with the suggested model, the diffusion of carbon atoms present in CNTs structures act as the predominant mechanism for vertical growth of Al4 C3 [131]. Figure 16 gives a schematic illustration of vertical and parallel growth of the carbides on CNTs.

Figure 16. A schematic view of growth mechanisms of Al4 C3 dispersoids on CNTs. It indicate the vertical migration of Al and C atoms for the formation of Al4 C3 carbides [131] (Reproduced with permission from [131], Elsevier, 2009).

As a summary, the formation, volume fraction, morphology, and size of the formed carbides in CNT-Al matrix nanocomposites depend on the processing parameters such as temperature, post treatment and purity of the initial CNTs. The in situ formed intermetallics can improve the mechanical properties of CNT-reinforced Al matrix nanocomposites through enhanced load transfer between the matrix and the reinforcements. (b) CNT-Reinforced Ti Matrix Nanocomposites A review on the literature confirms that TiC is the most common in situ formed intermetallic in CNT-reinforced Ti matrix nanocomposites. This compound often forms at the boundaries of the sintered powders, if the initial CNTs are wrapped by initial Ti powders. The minimum reaction temperature required for the generation of TiC precipitates is 800 ◦ C. The evaluation of the kinetic aspects of TiC formation in CNT-Ti binary systems has proved the faster formation kinetics of TiC compared to CNTs. The volume fraction of the formed TiC dispersoids differs with the chemical composition of the incorporated carbon allotrope into Ti matrix, e.g., CNT and graphite. TiC particles may form in two different morphologies including spherical and elongated ones. The applied

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dispersion method for mixing the initial powders, processing parameters such as sintering temperature and the initial volume fraction of CNT, as well as the application of post heat treatments can remarkably affect the mechanical properties of the composites. Generally, the empirical results have demonstrated the positive effects of the formed intermetallics on the physicomechanical properties of the nanocomposites. There are several intermetallic compounds which can be formed during the reaction between Ti and carbon allotropes. Among these compounds, TiC is the most common [18,140]. It can be formed during a solid-state reaction between Ti and CNTs [18,140,141]. The standard free energy of the reaction may be obtained using the following equation: ΔG = −184,571.8 + 41.382T − 5.042T × lnT + 2.425 × 10−3 T2 − 9.79 × 105/T (T < 1939 K)

(3)

The minimum reaction temperature required for the generation of TiC precipitates in CNT-Ti binary systems is 800 ◦ C. These particles often locate at the boundaries of the sintered powders, provided that the initial CNTs are wrapped by initial Ti powder. It is found that reaction between the nanotubes and Ti matrix is promoted by an increment in the sintering temperature [141]. As to CNTs, the mentioned reaction is fast and produces more TiC particles. Thus, TiC dispersoids are the dominant phase in the microstructure. The volume fraction of the formed TiC precipitates strongly depends on which carbon allotrope is incorporated into Ti matrix. Figure 17 shows the microstructure of Ti matrix reinforced with 0.4 wt % CNTs or 0.4 wt % graphite. As clearly seen, higher amounts of TiC dispersoids are present throughout Ti matrix in the MWNTs-reinforced composite (as indicated by white arrows) [142].

Figure 17. Optical images of Ti matrix reinforced with: (a) CNTs and (b) graphite [142] (Adapted with permission from [142], Elsevier, 2013).

TiC particles can form in different morphologies. For instance, Kondoh et al. [140] have shown the formation of two different shapes of TiC particles, namely spherical particles and elongated ones. Spherical particles are evolved as a result of incomplete disassembling of MWNTs, while extrusion-induced severe plastic deformation leads to the generation of elongated ones. Thus, the post treatment can affect the morphology of the formed intermetallics. Figure 18 shows the microstructure of the extruded composites, wherein Figure 18a indicates the uniformly distributed reinforcements throughout Ti matrix and Figure 18b confirms the desirable bonding between Ti matrix and TiC particles. The experimental measurements confirm the positive influence of TiC precipitation at CNT/Ti interfaces on the physicomechanical properties of these nanocomposites [18,35,50,140]. The higher the volume fraction of incorporated CNTs, the superior the mechanical properties of the composites are [18,140]. However, the exact strengthening mechanism of intact CNTs has not been clarified [18]. As practical variables, the used method for mixing the initial powders, processing parameters such

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as sintering temperature, and initial volume fraction of CNTs, as well as the post heat treatment can remarkably affect the mechanical properties of the composites [18]. In the case of the mixing method, Munir et al. [143] have demonstrated that unlike the solution ball milling (SBM) process, TiC can be formed when the initial powders (Ti powder + 0.5 wt % MWNTs) are mixed by high energy ball milling. However, the in situ TiC particles can be formed by using both methods whenever 1.0 wt % MWNTs are incorporated into Ti matrix. As to the potential effects of the processing parameters on the mechanical behavior of CNT-Ti matrix nanocomposites, Xue et al. [141] have demonstrated that the compressive yield strength of the nanocomposites decreases with an increase in the sintering temperature before a slight rise. This is originated from the combined effect of increased relative density and structural damage of CNTs. Finally, the relative density of in situ formed TiC dispersoids increases with an increment in the volume fraction of initial CNTs. Thus, the higher the initial CNTs volume fraction, the better the mechanical properties will be [142]. In the case of the way a post heat treatment can affect the mechanical behavior, it is shown that the properties can be degraded upon increasing the post annealing temperature. This decrease is “gradual”, because the presence of TiC dispersoids throughout Ti matrix may drastically prevent Ti particles from the extreme coarsening through the particle-pinning mechanism [18].

Figure 18. Extruded CNT-dispersed metal matrix nanocomposites: (a) Optical image of the binary systems and (b) SEM image of TiC/Ti interface [140] (Reproduced with permission from [140], Elsevier, 2009).

As a summary, the incorporation of CNTs into the Ti matrix can drastically enhance the tensile strength, yield strength, and microhardness values of the matrix. It may be ascribed to the dispersion effect of un-bundled CNTs and in situ formed TiC particles, as shown in Figure 19.

Figure 19. SEM image of fracture surface of CNT-Ti nanocomposites in different magnifications [140]: (a) ×5000; (b) ×65,000 (Adapted with permission from [140], Elsevier, 2009).

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(c) CNT-Reinforced Mg Matrix Nanocomposites Mg and its alloys are not able to sufficiently wet CNTs. To solve this practical challenge, two general approaches have been suggested: (i) Coating of CNTs by a metallic thin film and (ii) alloying Mg matrix with appropriate elecments. The first solution can efficiently enhance the interfacial adhesion through the prevention of CNTs from clustering. Clustering of CNTs adversely alter the interfacial bonding, resulting in deteriorated mechanical properties. It is found that an intermetallic compound (Mg2 Ni) can be generated through wrapping CNTs by a Ni-based coating and alleviate decreased properties in Mg matrix composites. It is ascribed to an improved bonding at the interface, enhancing the mechanical properties of the nanocomposites with no significant decrement in ductility [12]. From this perspective, the utilization of Mg-Al alloys (AZ-type) such as AZ31, AZ81, Mg-3Al-1Zn, and Mg-6 wt % Al instead of pure Mg is strongly recommended because of the easy formation of ternary Al2 MgC2 compounds. Depending on the chemical composition of used Mg alloy matrix and CNTs coatings, the several intermetallic phases of different stoichiometries may form among which Al2 MgC2 , Al3 Mg2 , Al12 Mg17 , Al4 C3 , and Mg2 Ni are of prime significance [144]. Figure 20 shows the needle-like morphology of Al2 MgC2 ternary carbides.

Figure 20. TEM image of (Mg-6 wt % Al)/CNT interface, indicating the in situ formation of needle-like ternary carbides (Al2 MgC2 ) [144] (Reproduced with permission from [144], Elsevier, 2011).

The precipitation of the interfacial intermetallics has an optimum limit, so that the mechanical properties of the nanocomposites may be weakened whenever it is exceeded. Therefore, a close control over the interfacial reactions between CNTs and Mg is highly required. These reactions proceed by simultaneous diffusion of the matrix atoms and carbon atoms present on the superficial regions of CNTs toward the preferred reaction sites. To confine Mg-CNTs reactions along the interfaces, employing a short-term method such as microwave sintering has been suggested. The long-time contact between the matrix and reinforcements may drastically degrade the mechanical properties of the nanocomposites due to the formation of excessive content of the intermetallics [12]. As previously discussed, if the precipitation rate of the intermetallic compounds are closely controlled, enhanced properties can be obtained. This enhancement can be justified by three different mechanisms: (i) Grain refinement, (ii) improvement in load transfer from the matrix to reinforcements, and (iii) appropriate wettability. For instance, it is shown that the Al4 C3 phase often acts as a suitable grain refiner and boosts the tensile strength of CNT-AZ81 nanocomposites [19]. As another example, the formation of SiC during the sintering of MWNTs-AZ91 powder blend facilitates the load transfer from the matrix to MWNTs and enhances the mechanical properties of produced composites [145]. In addition, the mechanically improved behavior of (Ni-coated MWNTs)-Mg composites is ascribed to

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the presence of an adherent and void-free MWNTs /Mg interface due to Mg2 Ni precipitation. Figure 21 shows Mg2 Ni dispersoids at the MWNTs/Mg interfaces [12].

Figure 21. TEM image of (Ni-coated MWNT)-Mg composites, showing the presence of Mg2 Ni intermetallic at the interface [12] (Adapted with permission from [12], Elsevier, 2014).

Similar to CNT-Al and CNT-Ti binary systems, the precipitation of the mentioned intermetallics may bilaterally affect the mechanical properties of CNT-dispersed Mg matrix nanocomposites. The poor failure strain of the composites is attributed to the ubiquitous presence of coarse intermetallic particles. The initial volume fraction of incorporated CNTs, formation of intermetallics, and quality of the formed interface between the Mg and nanotubes are the important factors in determining the mechanical properties of the nanocomposites. It is shown that an increase in the initial amount of incorporated CNTs may promote the refinement of coarse grains and enhance the failure strain [146,147]. Figure 22 shows the evolved microstructure of CNT-AZ31 nanocomposites produced using the disintegrated melt deposition technique followed by hot extrusion. As seen, the formed Mg17 Al12 particles are distributed at or along α phase boundaries and give rise to the grain refinement of Mg matrix [148].

Figure 22. Optical image of (a) AZ31 and (b) 1 vol % CNT-AZ31 nanocomposites fabricated through the disintegrated melt deposition technique followed by hot extrusion at 350 ◦ C. In these systems, Mg17 Al12 intermetallics are abundantly formed. These compounds are manifested by black regions, while white regions exhibit α-Mg [148] (Reproduced with permission from [148], Springer, 2012).

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The formation of the mentioned intermetallics enhances the mechanical properties of CNT-Ti matrix nanocomposites through guaranteeing the suitable wetting between the matrix and nanotubes. It can significantly prevent the microscale cavities which form at the Ti/CNT interfaces and contribute to improved mechanical properties [12]. Additionally, the formed intermetallics can improve the mechanical properties of the nanocomposites through the restriction of mechanical deformation [148]. The formation of a clean interface with no impurity or defect is shown to improve the mechanical properties [144]. A recent study by Rashad et al. [146] showed that the formed intermetallics in MWNTs-dispersed Mg-3Al-1Zn alloys are refined with the CNTs inclusion. Figure 23 indicates the schematic demonstration of Mg17 Al12 generation at the interface of Mg-3Al-1Zn/MWNTs composite, where the coarse particles of Mg17 Al12 can be refined through the incorporation of MWNTs into the matrix.

Figure 23. A schematic view of Mg17 Al12 particles formed in Mg-3Al-1Zn/MWNTs nanocomposites prepared by powder metallurgy method. The incorporation of MWNTs can drastically refine the particles (β phase corresponds to Mg17 Al12 ) [146] (Reproduced with permission from [146], Elsevier, 2015).

A review on the literature confirms that recent research works have strived to attribute the enhanced properties of CNT-Mg composites to the formation of different intermetallics. Fukuda et al. [144] have resorted to Al2 MgC2 compound. However, an opposite trend has been observed by Pei et al. [110], where Al2 MgC2 strengthens the carbon nanofiber/Mg interface, but simultaneously leads to decreased mechanical properties. It is ascribed to the difference in the extent of graphitization. In other words, produced Al2 MgC2 intermetallic in CNT-reinforced composites insignificantly affects the CNTs strength. Table 4 provides the readers with an overview of the fabrication method, chemical composition of the possible intermetallic compounds, and variations in mechanical properties of the composites.

29

30

3 vol % MWNTs

1.0 vol %

DMD method followed by hot extrusion

Accumulative roll bonding

1–5 wt % MWNTs

Squeeze infiltration method

0.25–1.0 wt % MWNTs

1.5 vol % CNT nanoparticles

DMD (disintegrated melt deposition) technique followed by hot extrusion

Powder metallurgy method

0.3 wt % MWNTs

Powder metallurgy technique

Unfunctionalized

Unfunctionalized

Acid treated

Acid treated

Unfunctionalized

Al3 Mg2 , Mg17 Al12

Mg17 Al12

Mg17 Al12

SiC

Al4 C3 , Al12 Mg17 , Al3 Mg2

Mg2 Ni

Al2 MgC2

Functionalized; using a zwitterionic surfactant

0.71–1.56 vol % MWNTs

Powder metallurgy based wet processing

Ni-coated

Chemical Composition of the Formed Intermetallics

Type of Used CNTs

CNTs Volume Fraction

Fabrication Method

An increase in microhardness, yield strength, and ultimate strength by 16 HV, 80 MPa, and 66 MPa, respectively. A decrease in ductility by 1.5%

-

An increase inmicrohardness, yield strength, ultimate strength, and elongation at failure by 29 HV, 34 MPa, 8 MPa, and 3.6%, respectively An increase in ultimate strength and failure strain by 10 MPa and 9%, respectively Yield strength remains unchanged

-

-

-

An increase in microhardness, yield strength, and ultimate strength by 80 HV, 170 MPa and 90 MPa, respectively

-

-

An increase in yield strength, ultimate strength, and failure strain by 107 MPa, 114 MPa, and 7%, respectively

-

-

An increase in yield strength and elongation at failure by 28.5 MPa and 12.2%, respectively

-

Variations of Mechanical Properties

Grain refinement due to the presence of Al4 C3 Crack formation-based tensile toughening mechanism Grain refinement Improved load transfer owing to SiC formation at the interface Grain refinement Strengthening by stable intermetallic phases CNTs-induced strengthening

-

-

CNTs-induced strengthening Refinement of coarse Mg17 Al12 particles through the incorporation of CNTs into the matrix, which improves the ductility CNTs-induced strengthening Refinement of coarse Mg17 Al12 particles through incorporation of CNTs into the matrix, which improves the ductility

-

-

-

-

Grain refinement CNTs-induced strengthening Improved interfacial interactions

Presence of clean interface CNTs-induced strengthening Formation of Al2 MgC2

-

-

Suggested Reasons for Enhanced/Deteriorated Properties

Table 4. An overview of fabrication method, formed intermetallic compounds, and mechanical properties of CNT-reinforced Mg matrix nanocomposites.

[147]

[146]

[148]

[145]

[19]

[12]

[144]

Ref.

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As another interesting result, a slight increment in relative density of CNT-Mg nanocomposites has been observed owing to the formation of in situ intermetallics. For instance, it is found that the incorporation of 0.3 wt % Ni-coated CNTs into Mg matrix increases its density by 0.01 g/cm3 [12]. (d) CNT-Reinforced Cu Matrix Nanocomposites Similar to Mg and its alloys, Cu is not able to sufficiently wet CNTs. This is because there is a significant difference between the surface tension of Cu and CNTs. Furthermore, it is not expected that Cu reacts with CNTs in conformity with Cu-C phase diagram. Instead, some fissures may form at Cu/CNTs interfaces [149]. Figure 24a shows such an obvious crevice at the interface. The addition of Ti to Cu matrix eliminates these crevices through the formation of a thin TiC layer at the interface, as shown in Figure 24b. Such crevices can significantly degrade the final properties of the nanocomposites, e.g., thermal conductivity, which will be discussed in the next section.

Figure 24. TEM images of CNT/Cu interfaces: (a) The formation of a pronounced crevice due to the lack of wettability and (b) the precipitation of in situ TiC layer as a result of Ti addition to pure Cu matrix [149] (Reproduced with permission from [149], Springer, 2013).

The improvement of the interfacial bonding between Cu matrix and CNTs has been the focus of many research works. One of the suggested solutions is alloying of Cu matrix using carbide forming elements such as Cr and Ti. It can lead to the precipitation of Cr3 C2 , TiC, and Cr7 C3 carbides. The open ends of CNTs and carbon atoms present on the wall defects can serve as the carbon source reacting with Cr [150]. These carbides are generated at the appropriate sites including broken surfaces and reactive edges of the tubes through various bonding mechanisms such as rope anchors, chain-link structures and bridges [151]. The standard free energy (ΔGTθ ) of the reactions may be obtained using the following equations [152]: (4) 3Cr + 2C(g) → Cr3 C2 ΔGTθ = − 72.333T × ln T + 0.062T 2 + 1, 345, 000T −1 − 0.00001T 3 + 744.180T − 110, 762.634

7Cr + 3C(g) → Cr7 C3

(5)

ΔGTθ = −114.064T × ln T + 0.108T 2 + 1, 772, 000T −1 − 0.0000192T 3 + 706.278T − 218, 451.845

Ti + C → TiC ΔGTθ

= −27.75T × ln T +

0.0241T 2

+ 665, 300T

−1

−

0.0000003T 3

(6) + 191.888T − 194, 857.227

In general, Cu-Ti compounds precipitate during SPS, if the Cu-Ti alloy is used as matrix. It is while the Cu-Ti compounds may be decomposed into TiC during the sintering process. The formation of TiC proceeds through the reaction between CNTs and Ti diffused from the primary particle boundaries [20].

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An island-shaped Cr7 C3 can form as a result of a chemical reaction between the open tip of MWNTs and diffused Cr. Other carbides can be evolved due to a reaction between defective sidewalls of MWNTs and Cr atoms [49]. It is shown that the formed carbides at the tip of MWNTs can provide superior load transfer, because they are in contact with more amounts of graphene layers [153]. The studies show that the controlled formation of C-Cu intermetallic compounds at Cu/CNTs interfaces can noticeably improve the physicomechanical properties, thermal conductivity, pitting corrosion resistance, and electrical conductivity of these binary composites. The major strengthening mechanism responsible for this enhancement is the improved load transfer from the matrix to CNTs as a result of carbides precipitation. From mechanical point of view, the Cr7 C3 particles are reported to effectively transfer the applied tensile loads to MWNTs in CNT-(Cu-Cr) nanocomposites [153]. Also, the yield strength of Cu-Ti alloy increases by 88% through the inclusion of CNTs. It is ascribed to the synergic combination of plastic deformation and stronger interfacial bonding induced by the formation of a thin TiC layer [154]. The literature confirms the degraded thermal conductivity of MWNTs-reinforced pure Cu matrix composites due to the poor interfacial bonding at the CNT/Cu interface. Alloying of Cu matrix may modify superficial regions of MWNTs and drastically enhance the thermal conductivity of the nanocomposites due to increased heat transfer. Improved interfacial bonding is attributed to the carbides formation. Another suggested approach to improving the interfacial bonding is the coating of MWNTs by a metallic layer. Nonetheless, the thermal conductivity of these composites may be considerably degraded due to the formation of excessive impurities during the coating process. In contrast, the generation of metallurgical bonds between Cu-Ti alloy matrix and MWNTs improves the thermal conductivity of the composites due to the formation of TiC dispersoids. This compound can facilitate the electron-phonon coupling and prevent their scattering [149]. As another typical example, the superior thermal conductivity of CNT-(Cu-Cr) nanocomposites can be achieved through the generation of Cr3 C2 at the interface. These precipitates are perfectly capable of improving the heat transfer efficiency by strengthened interfacial bonding. Figure 25 indicates the precipitation of Cr3 C2 at the CNT-(Cu-Cr) interface. As a general rule, if the intermetallic compounds form at the interface of these composites, the thermal conductivity will be enhanced. Furthermore, the volume fraction of initial CNTs strongly affects the thermal conductivity of the nanocomposites, wherein the thermal conductivity increases with an increment in the initial volume fraction of CNTs, as shown in Figure 26.

Figure 25. TEM image of Cr3 C2 precipitates at the CNT-(Cu-Cr) interface [150] (Adapted with permission from [150], Springer, 2013).

As to the electrochemical activity, the presence of the mentioned intermetallics may increase the pitting corrosion resistance of MWNTs-dispersed Cu matrix nanocomposites. This effect is related to enhanced passivation originated from the formation of the carbides. This is the case for the electrical conductivity. The enhancement in electrical properties of Cu-Ti solid solutions reinforced with MWNTs 32

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can be obtained upon the formation of ubiquitous intermetallic compounds. These carbides diminish the concentration of Ti atoms in the matrix solid solution and statistically reduce the scattering sites of charge carriers [20]. For the pure Cu matrix, a similar improvement is observed. It is ascribed to enhanced electron transport as a result of strong interfacial bonding originated from TiC formation at MWNTs-Cu interfaces [154].

Figure 26. Thermal conductivity of CNT-(Cu-Cr) nanocomposites as a function of initial volume fraction of CNTs [150] (Reproduced with permission from [150], Springer, 2013).

In summary, the mechanical, corrosion-related, and electrical properties of the nanocomposites can be enhanced through the in situ formation of some intermetallics as a consequence of improved load transfer, enhanced passivation and suitable electron transport. (e) Other Composites Systems Intermetallic compounds can be formed at the interface of pure Ni matrix and CNTs. In most cases, the formation of intermetallic compounds such as a metastable hexagonal Ni3 C at the defect sites of tubes can significantly improve the mechanical properties of CNT-dispersed Ni matrix nanocomposites due to the formation of strong interfacial bonds [52]. Such a scenario is observed in Fe alloys. In these alloys, the iron carbides form as a result of a chemical reaction between the iron matrix and included CNTs. Microstructure-related and mechanical properties of the composites are drastically affected by the carbides formation. Lin et al. [155] proved the deterioration of mechanical properties in these composites with carbides precipitation; albeit the reason has not been discussed. Partial Chemical Reaction The strength of CNT/metal interfacial bonding is one of the key factors controlling the physicomechanical features of CNT-reinforced metal matrix nanocomposites as well as the uniform distribution of CNTs and high relative density [49]. As a general rule, the weak wettability and poor interfacial compatibility between CNTs and metal matrix degrades the fracture toughness due to the limited interfacial load transfer [156]. Among the widely used methods for the fabrication of metal matrix nanocomposites (MMNCs), powder metallurgy routes provide a better interfacial bonding as compared to casting processes [157]. A model is developed by Coleman et al. [158] describing the tensile strength of CNT-containing composites (σc ) based on the shear strength of the interface (σshear ). Equation (7) describes this model:  σC =

1 +

b R



lCNT − σ 2R shear

 1 +

  b σM VCNT +σM R

(7)

where b is thickness of the interface layer, σM is the strength of the matrix and R, lCNT and V CNT are the radius, length and volume fraction of CNTs, respectively. 33

Metals 2017, 7, 384

Zhou et al. [159] used an in-situ pull-out technique to quantitatively evaluate interfacial shear strength (IFSS) in CNT-Al nanocomposites. They bonded the end of the protruding CNTs directly to the tip of an atomic force microscopy (AFM) cantilever and pulled them out from Al matrix. They reported that the estimated tensile strength of the nanocomposite based on the shear lag model using the obtained IFSS values is consistent with the experimental ones. To date, the direct microscopic method for determination of IFSS values have also been employed for CNT-filled ceramic [160] and polymer [161] matrix nanocomposites. It is inferred that the tensile strength of the nanocomposites can be maximized by thickening the interfacial region through the formation of a crystalline coating around CNTs or their functionalization or in situ formation of metal oxide/carbide interlayer at boundary zones. For instance, shear strength of CNT-Al nanocomposites is enhanced by the formation of interfacial Al4 C3 carbides during a chemical reaction between defective CNTs and molten Al in the interboundary areas [120,162]. Figure 27 shows the interfacial regions in CNT-Al matrix nanocomposites. As seen, several in situ phases are embedded in the grain boundary layer of Al-CNT sintered compacts including Al2 O3 , amorphous carbon black, graphite and Al4 C3 phases [120]. The formation of carbides on the surface of CNTs at Al/CNT interfaces diminishes the contact area between CNTs and Al matrix with poor wettability and substitutes the Al/Al carbide interfaces with a lower wetting angle than that of CNT/Al boundaries (~55◦ vs. 130–140◦ ) [162,163]. The stronger the physicochemical interactions between these two phases, the more adhesive the interface will be. To simply describe metal/ceramic physical interaction, the work of adhesion (W ad ) is defined as: (i) The amount of energy released whenever two free surfaces are brought into contact or (ii) the work per unit area of interface which is reversibly performed to separate the two phases. W ad can be given by the following equation [128,164]: Wad = σLV (1 + cos θ )

(8)

where σLV is the surface tension and θ is the contact angle between two phases. Therefore, W ad as a measure of the interfacial adhesion could be enhanced by improving the wettability through extensive interfacial chemical reactions. From this point of view, Al4 C3 formation at the CNT/Al interface increases the interfacial adhesion from 200 mJ/m2 for Al-C system to1156 mJ/m2 for Al-Al4 C3 (calculated at 1100 ◦ C) and consequently counteracts delamination and improves the load transfer efficiency [128].

Figure 27. TEM images of the grain boundary regions in CNT-dispersed Al matrix nanocomposites: (a) Micrograph of the whole grain boundary including (1) Al, (2) alumina, (3) CNT, (4) amorphous carbon black, (5) graphite, and (6) Al4 C3 phases; (b) interfacial region between CNT and Al matrix; and (c) HRTEM image and SAD (Selected area diffraction) pattern of Al4 C3 phase. The region with the broken alumina phase is indicated with the white arrow [120] (Reproduced with permission from [120], Elsevier, 2009).

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On the contrary, some research works have reported the adverse effects of the carbide formation on the CNT/Al interfacial bonding and degradation of mechanical properties in CNT-Al nanocomposites due to a weak interfacial adhesion [165–167]. However, some researchers have demonstrated directly or indirectly the desired influences of CNT/metal matrix interfacial reactions on the mechanical features of the nanocomposites regarding the fact that the CNT/metal interfacial load transfer is a direct function of the interfacial bonding. Zhou et al. [168] produced the Al4 C3 nanostructure at the end of CNTs incorporated into Al matrix through an appropriate heat treatment and investigated the thermal expansion behavior of the nanocomposite as a criterion of CNT/Al matrix interfacial bonding. They observed stress contrast around carbides as an evidence of a friction trace, leading to the enhancement in their anchor effect from Al matrix. This carbide-induced anchor effect minimizes the local interfacial slippage and constrains the matrix deformation. Consequently, the thermal expansion exhibits a linear and reversible behavior under cyclic thermal load, indicating a reduced interfacial slippage between CNTs and Al matrix at the presence of Al4 C3 precipitates. It is proven that CNT/metal interfacial adhesion strongly depends on the applied method for dispersion of CNTs inside the matrix. Among the various discussed methods, molecular-level process and metallization of CNTs are the most effective. Kim et al. [49] used molecular-level process prior to spark plasma sintering to fabricate CNT-filled Cu nanocomposites. They reported enhanced hardness and wear resistance in spark plasma sintered CNT-Cu nanocomposites than pure Cu with good CNT/Cu interfacial bonding as well as uniform CNTs distribution in the matrix and high relative density. Metallization of CNTs using chemical vapor deposition (CVD) [16] and electroless coating [95] is one of the effective approaches to achieve a strong metal/CNT interfacial bonding as well as a homogeneous dispersion of CNTs inside the matrix. It can improve the load transfer from matrix to CNTs through the metallic coating. In another study, He et al. [128] fabricated CNT-reinforced Al nanocomposites with enhanced mechanical properties via the in situ chemical vapor deposition process and ascribed the strong interfacial bonding to the formation of transition thin layer of Al4 C3 between CNTs and Al matrix. Clean oxide-free surfaces of metallic powders are a prerequisite for obtaining a good CNT/metal interfacial adhesion. To meet this challenge, some additives in CNT-metal systems are utilized to reduce the oxide content. Owing to its negligible solid solubility, good adhesion with Cu, and high thermodynamic stability, 0.5 wt % ruthenium (Ru) as an air-stable transition metal was exploited in a recent survey to reduce non-protective oxide films on Cu surfaces. Furthermore, Ru contributes to increased thermal conductivity of CNT-Cu binary systems as a result of enhanced CNT/Cu interfacial adhesion and inhibition of the heat carrying phonons scattering at the CNT/Cu interfaces and weak interfacial bonding-induced pores. However, if Ru content exceeds 2.5 wt %, a reduction in hardness is observed due to the restricted densification of the powder system. It arises from high melting point of Ru, serving as a diffusion barrier during the consolidation process. Also, using 1 wt % Ru decreases the electrical conductivity of pristine copper, because the interfacial resistance occurs due to high electrical resistivity of Ru than Cu (7.1 × 10−6 vs. 1.68 × 10−6 Ω·cm) [169]. Interestingly, metal oxide formation on the surface of the particles may strengthen the interfacial adhesion in some CNT-reinforced metallic systems. For instance, Kondoh et al. [97] fabricated CNT-Mg nanocomposites by spark plasma sintering of CNT-coated Mg particles and reported in-situ formation of MgO dispersoids during the consolidation process as well as a thin MgO film on the surface of Mg particles due to the atmospheric corrosion. They confirmed that the diffusion of carbon atoms into the MgO structure provides an enhanced interfacial adhesion with CNTs. This phenomenon is mechanistically similar to the epitaxial growth of chromium carbide nanostructures mostly Cr7 C3 at the CNT/Cu-Cr interface as a result of the substitutional diffusion of Cr atoms into radially unzipped defects of CNTs. It is schematically shown in Figure 28. The results show that the formation of Cr carbide nanostructures may potentially increase the efficiency of interfacial load transfer while preserving CNTs structure [170]. It is further surveyed by Kathrein et al. [171].

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Figure 28. A schematic view of epitaxial growth of Cr carbide at the CNT/Cu-Cr interface via the substitutional diffusion of Cr atoms into the nanotubes during sintering [170] (Reproduced with permission from [170], Elsevier, 2013).

The nanocomposite processing parameters such as applied temperature/pressure and post thermal/mechanical treatments are among factors affecting CNT/matrix interfacial adhesion. Guo et al. [172] fabricated CNT-Al nanocomposites through spark plasma sintering followed by hot rolling process and reported that higher sintering temperature (630 ◦ C vs. 590 ◦ C) provides a stronger cohesion force for Al-CNTs and Al-Al particles due to higher degree of densification. 3. CNT Pinning-Induced Grain Refinement In addition to the load bearing capacity, CNTs can affect the grain size of the metal matrix nanocomposites. In fact, the incorporation of CNTs into a metallic matrix results in lower grain size due to the CNT-induced pinning effect. Such an effect arises from the low dimensions of CNTs which can pin the grains of the metal matrix similar to nanosized round particles. The more the CNT content, the higher the effectiveness of boundary pinning and the lower the grain size are [42,101,173,174]. Figure 29 shows the STEM (Scanning transmission electron microscope) micrographs of Cu and CNT-Cu nanocomposites after the consolidation by high pressure torsion (HPT). As obviously seen, an increase in CNT content may feasibly reduce the grain size and narrow its distribution thank to the grain pinning effect [42]. The mean grain size (D) of reinforcement-containing metal matrices can be estimated by the Zener pinning relation [174]: k×r D= (9) fn where k is a proportional dimensionless constant, f is the volume fraction of the secondary phase, and r is the mean reinforcement radius, respectively. The secondary phase particle (i.e., CNTs) acts as a frictional force against the grain boundary migration and hinders the grain growth. As seen in Equation (7), the strengthening effect of particles depends on their radius (diameter). The smaller the diameter of CNTs, the smaller the mean grain size of the metallic matrix will be. Moreover, the average distance between CNTs is another factor which greatly affects the pinning effect. An increase in CNT content gives rise to lower distance between CNTs, increased interfacial energy absorption and decreased grain size. Additionally, for a constant CNT content, the average distance between short CNTs are lower than that between long ones. As a result, short CNTs have stronger pinning effect and result in lower grain size [174]. Such a grain refinement affects the mechanical properties of CNT-dispersed metal matrix nanocomposites. In general, CNTs can increase the mechanical strength of metal matrices through a variety of mechanisms: (i) Load transfer from the metallic matrix to CNTs, (ii) reduction in the matrix grain size through the pinning effect (Hall-Petch relation), (iii) CNT-induced Orowan looping mechanism, (iv) solid solution strengthening due to the diffusion of carbon atoms from CNTs to

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the matrix, (v) secondary phase (particle) strengthening due to the in-situ formation of carbide particles by a chemical reaction between CNTs and the metal matrix, (vi) work hardening of the matrix due to the formation and accumulation of dislocations at CNT/metal interfaces caused by thermal mismatch between CNTs and metal matrix, and (vii) the strengthening effect induced by impurities originating from the mixing methods [175]. Depending on the type of matrix and the fabrication method, one or some of these mechanisms may be stimulated. Another factor affecting the strengthening mechanisms is the physical and size-dependent properties of CNTs. For instance, if the thermal expansion coefficient of CNTs is considerably different with that of the matrix, the dislocations will build up at CNT/metal interfaces and enhance the mechanical strength of the nanocomposite. Moreover, the diameter and length (i.e., aspect ratio) of CNTs may influence the strengthening mechanisms in metal matrices. Among the aforementioned strengthening mechanisms, the load transfer mechanism is the most influential. Typically, the higher length or lower diameter of CNTs (i.e., higher aspect ratio) results in better load transfer and higher mechanical properties such as improved yield strength and elastic modulus. However, when the aspect ratio of CNTs is lower than a critical value, the Orowan mechanism is dominant [175,176]. Beside the positive effect of long CNTs on enhanced mechanical properties than short ones, long nanotubes are usually more sensitive to agglomeration. As a consequence, the dispersion of shorter CNTs is an easier task than that of longer ones [34]. Using a high content of long CNTs may induce the agglomeration and degrade the mechanical properties of final nanocomposite.

Figure 29. STEM (scanning transmission electron microscope) images of Cu and CNT-Cu nanocomposites fabricated by high pressure torsion (HPT): (a) Pure Cu, (b) 5 vol % CNT-Cu, and (c) 10 vol % CNT-Cu nanocomposites. The left images are provided by bright-field mode, and the right ones by high-angle annular dark-field mode [41] (Reproduced with permission from [41], Springer, 2013).

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4. Conclusions and Outlook The present review paper has provided a broad overview of practical challenges in the fabrication of CNT-dispersed metal matrix nanocomposites. These challenges have been categorized into four main groups: (i) Non-uniform dispersion of CNTs throughout the metallic matrix, (ii) thermal decomposition of CNTs and chemical reaction with the metallic matrix, (iii) poor interfacial adhesion, and (iv) low compactability. One can obtain a CNT-metal matrix nanocomposite with superior properties if all of four challenges are appropriately overcome. A large number of research works have focused on these issues and strived to suggest feasible solutions to them. Their strategy is often based on the better dispersion of CNTs through the effective prevention of their agglomeration and close control over the processing parameters for suppressing the unwanted phase transformations and preventing from unfavorable microstructural features. Although these solutions remarkably enhance the physicomechanical properties, each of them has its own limitations. It seems that the future research works should find their way toward development of new solutions for uniform dispersion of CNTs in metallic matrices, so that the electrical and thermal properties are improved as well as mechanical ones. It seems that the future research works should focus on the development of more efficient ways to uniformly disperse CNTs in the metallic systems or introduction of new methods for efficient consolidation of metallic powders mixed with CNTs. Acknowledgments: The present work has been conducted by some active members of the “Advanced Materials Research Group (AMRG)” founded by Abolfazl Azarniya in 2016. The authors would like to acknowledge all team members including Anne Jung (from Saarland University, Saarbrücken, Germany), Mohammad Mirzaali (from Delft University of Technology, Delft, The Netherlands), Dariusz Garbiec (from Metal Forming Institute, Poznan, ´ Poland), Ridvan Yamanoglou (from Kocaeli University, Kocaeli, Turkey), Temel Varol (from Karadeniz Technical University, Trabzon, Turkey), Mohammad Abedi and Dmitry Moskovskikh (from National University of Science and Technology MISIS, Moscow, Russia), Flávio Bartolomeu and Georgina Miranda (from University of Minho, Braga, Portugal), Xabier Garmendia (from University of Liverpool, Liverpool, UK), Joseph Ahn (from Imperial College London, London, UK), Yaser Shanjani (from Stanford University, Stanford, CA, USA), Chor Yen Yap, Sing Swee Leong, and Wai Yee Yeong (from Nanyang Technological University, Singapore), Marek Weglowski (from Institute of Welding, Gliwice, Poland), Wessel Wits (from University of Twente, Enschede, The Netherlands) and some others for their best efforts. These scientific partners are a part of AMRG’s forthcoming research publications. Author Contributions: Abolfazl Azarniya, Mir Saman Safavi, Saeed Sovizi, Amir Azarniya, and Biao Chen outlined the review, performed the literature search and wrote the manuscript. Hamid Reza Madaah Hosseini and Seeram Ramakrishna helped in compiling the wide information related to the topic and data presentation and performed technical editing for all corrections. Hamid Reza Madaah Hosseini and Seeram Ramakrishna also helped in finalizing the manuscript and developed the idea of the topic and were responsible for the correspondence. Conflicts of Interest: The authors declare no conflict of interest.

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metals Article

Enhancing the Hardness and Compressive Response of Magnesium Using Complex Composition Alloy Reinforcement Khin Sandar Tun 1 , Yuming Zhang 2 , Gururaj Parande 1 , Vyasaraj Manakari 1 and Manoj Gupta 1, * 1 2

*

Department of Mechanical Engineering, National University of Singapore, Singapore 117576, Singapore; [email protected] (K.S.T.); [email protected] (G.P.); [email protected] (V.M.) Department of Mechanical & Industrial Engineering, University of Toronto St. George Campus, Toronto, ON M5S 1A1, Canada; [email protected] Correspondence: [email protected]; Tel.: +65-6516-6358

Received: 22 March 2018; Accepted: 14 April 2018; Published: 17 April 2018

Abstract: The present study reports the development of new magnesium composites containing complex composition alloy (CCA) particles. Materials were synthesized using a powder metallurgy route incorporating hybrid microwave sintering and hot extrusion. The presence and variation in the amount of ball-milled CCA particles (2.5 wt %, 5 wt %, and 7.5 wt %) in a magnesium matrix and their effect on the microstructure and mechanical properties of Mg-CCA composites were investigated. The use of CCA particle reinforcement effectively led to a significant matrix grain refinement. Uniformly distributed CCA particles were observed in the microstructure of the composites. The refined microstructure coupled with the intrinsically high hardness of CCA particles (406 HV) contributed to the superior mechanical properties of the Mg-CCA composites. A microhardness of 80 HV was achieved in a Mg-7.5HEA (high entropy alloy) composite, which is 1.7 times higher than that of pure Mg. A significant improvement in compressive yield strength (63%) and ultimate compressive strength (79%) in the Mg-7.5CCA composite was achieved when compared to that of pure Mg while maintaining the same ductility level. When compared to ball-milled amorphous particle-reinforced and ceramic-particle-reinforced Mg composites, higher yield and compressive strengths in Mg-CCA composites were achieved at a similar ductility level. Keywords: magnesium; high entropy alloy; composite; hardness; compressive properties

1. Introduction Magnesium (Mg) is the lightest of all structural metals and possesses the highest strength-to -density ratio. In addition, magnesium has other favorable advantages, including a high damping capacity, high dimensional stability, good machinability, good electromagnetic shielding characteristics, and recyclability. Accordingly, magnesium is the designers’ choice for possible production of lightweight vehicles to meet the demand of reducing greenhouse emissions [1–3]. Magnesium is mostly used in the form of alloys in commercial applications [4,5]. With the advent of composite technology, research interest has been placed on the development of high-performance magnesium composites. By a careful selection of matrix and reinforcing phases, newly formed composite materials with significant improvements in elastic modulus, strength, ductility, and coefficient of thermal expansion can be fabricated. Composite materials are attractive because they offer the possibility for combining useful engineering properties of individual elements, which is otherwise not possible from monolithic materials. The attractive physical and mechanical properties obtained from metal matrix composites (MMCs) have made them potential candidates for aerospace, automotive, and other structural applications [6]. For the fabrication of magnesium composites, various types of ceramic

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reinforcements, such as alumina (Al2 O3 ), yttria (Y2 O3 ), zirconia (ZrO2 ), silicon carbide (SiC), boron carbide (B4 C), and titanium carbide (TiC), have been used in Mg matrix [7–11]. In addition to ceramic reinforcements, research efforts have been made to synthesize Mg composites using metal particle reinforcements, such as copper, nickel, titanium, molybdenum, and aluminum [11,12]. Investigations have also been made on Mg composites containing hybrid reinforcements in the form of ceramic plus metal, ceramic plus ceramic, and ceramic plus carbon nanotube (CNT) besides single ceramic, metal, and CNT reinforcements [11,13–16]. In a recent development on magnesium composites, ball-milled amorphous particles were also used as a reinforcement in a pure magnesium matrix [17–19]. In the present study, an attempt is made to develop new magnesium composites using ball-milled complex composition alloy (CCA) particles. From the thermodynamic calculation, the current complex composition alloy has a high entropy value (ΔSmix ) of 12.97 J/K mol. Based on the concept of configurational entropy, the multicomponent alloys having a mixing or configurational entropy value which is equivalent to or greater than 12.471 J/K mol (ΔSconf ≥ 1.5R (12.471), where R is the gas constant), are regarded as high entropy alloys (HEAs). According to this concept, the CCA used in this study can be classified under the category of high entropy alloys. A careful examination of the published literature reveals that no attempt has been made to investigate the microstructure and mechanical properties of magnesium using ball-milled CCA particles. A powder metallurgy route incorporating microwave-assisted rapid sintering coupled with hot extrusion was used to synthesize Mg and Mg-CCA composites. The effect of the presence of CCA particles on microstructure and mechanical properties of magnesium was investigated. The interrelation between microstructure and mechanical properties of Mg-CCA composites was studied. The properties of Mg-CCA composites when varying the amount of reinforcement alloy particles are reported and the test results are benchmarked against pure magnesium. 2. Materials and Methods 2.1. Synthesis of Materials In this study, magnesium powder of 98.5% purity and with a size range of 60–300 μm (Merck, Darmstadt, Germany) was used as the matrix material. Ball-milled CCA particles were used as the particulate reinforcements. Initially, CCA pieces from cast ingot with composition Al35 Mg30 Si13 Zn10 Y7 Ca5 [20] were crushed into particles by ball milling the cast pieces at 200 rpm for 30 minutes with a ball to cast pieces weight ratio of 10:1 in a RETSCH PM-400 mechanical alloying machine (RETSCH, Haan, Germany). Monolithic magnesium and magnesium composites (Mg-2.5 wt % CCA, Mg-5 wt % CCA, and Mg-7.5 wt % CCA) were synthesized using a powder metallurgy technique. The synthesis process for Mg-CCA composites involved blending pure magnesium powder with CCA particles using the same machine at 200 rpm for 1 h without the use of grinding balls. The blended powder mixtures were then cold compacted using a 100-ton press to form billets that measured 35 mm in diameter and 45 mm in height. Monolithic magnesium was compacted using the same parameters without blending. The compacted billets were then sintered using a hybrid microwave sintering technique for 16 min to reach a temperature of 640 ◦ C near the melting point of magnesium using a 900 W, 2.45 GHz SHARP microwave oven (Sharp Corporation, Osaka, Japan). The sintered billets were homogenized at 400 ◦ C for 1 h and subsequently hot extruded at a temperature of 350 ◦ C at an extrusion ratio of 20.25:1. 2.2. Characterization Microstructural characterization studies were conducted on the extruded polished samples to determine the grain size, grain morphology, and presence and distribution of reinforcements. An Olympus metallographic optical microscope (Olympus Corporation, Shinjuku, Tokyo, Japan), MATLAB analysis software (R2013b, MathWorks, Natick, MA, USA), and a JEOL JSM-6010 scanning electron microscope (JEOL Ltd., Tokyo, Japan) were used for this purpose. X-ray diffraction analysis

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was conducted using an automated Shimadzu LAB-XRD-6000 (Shimadzu Corporation, Kyoto, Japan) (Cu Kα: λ = 1.54056 Å) spectrometer with a scan speed of 2 degrees per minute. Microhardness measurements were performed on the polished samples using a Shimadzu-HMV automatic digital microhardness tester (Shimadzu Corporation, Kyoto, Japan) with a Vickers indenter. An indentation load of 245.5 mN and a dwell time of 15 s was used in accordance with the ASTM standard E384-08. Room temperature compression tests were performed on cylindrical monolithic and composite samples according to ASTM E9-89a using an automated servo hydraulic testing machine MTS810 (MTS systems corporation, Eden Prairie, MN, USA). An extruded rod 8 mm in diameter was cut into 8 mm-long samples for compression tests to provide the aspect ratio (length/diameter) of unity. Samples were tested at a strain rate of 5 × 10−3 min−1 and the compression load was applied parallel to the extrusion direction. Fracture surface characterization studies were carried out on the compressively fractured surfaces of Mg and Mg composites with the objective of establishing the failure mechanisms. Fractography was accomplished using a JEOL JSM-6010 scanning electron microscope (SEM). 3. Results and Discussion 3.1. Analysis on Reinforcement Particles Figure 1a shows the size and morphology of the ball-milled CCA particles. The average particle size was calculated to be 2.7 ± 1.4 μm. The morphology of the particles can be seen as irregular shape although the smallest particles were almost spherical. The particle distribution can be seen in Figure 1b, and the particle sizes ranged from 1 μm to 7 μm with the dominant particle size in the range of 2–3 μm. The measured bulk microhardness of the cast materials was found to be very high at 406 ± 15 HV (4 ± 0.1 GPa) [20]. The XRD profiles of cast CCA and ball-milled CCA particles can be seen in Figure 2. The main phases found in the cast materials were maintained in the ball-milled version of the CCA material. The prominent change in the XRD pattern of ball-milled CCA particles is the broadening of the peaks, which may be attributed to the transformation of particles from the bulk material and an increase in the solid solubility of elements (Figure 2).

Figure 1. S scanning electron microscope (SEM) micrograph showing the morphology of ball-milled complex composition alloy (CCA) particles in (a) and the graph showing particle distribution in (b).

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Figure 2. X-ray diffraction (XRD) pattern of the cast CCA alloy and ball-milled CCA alloy particles.

3.2. Microstructure The distribution pattern of CCA reinforcement particles in the Mg matrix is shown in Figure 3b–d. For comparison purpose, the SEM micrograph of Mg is shown in Figure 3a. With the addition of 2.5 wt % CCA particles, the presence of uniformly distributed reinforcement particles can be seen in the micrograph (Figure 3b). In Mg-5CCA’s composition, the reinforcement particles were reasonably distributed in the Mg matrix (Figure 3c). However, with the increased addition of CCA particles from 2.5 wt % to 5 wt %, the tendency for the formation of particle clustering can be seen in the micrograph and the particles were mostly decorated at the particle/grain boundaries (Figure 3c). In the case of Mg-7.5CCA’s composition, in spite of having an increased amount of HEA particles, the HEA particles were uniformly distributed with limited evidence of particle clustering (Figure 3d) in the Mg matrix.

Figure 3. SEM micrographs of the: (a) Mg and CCA particle distribution in the (b) Mg-2.5CCA; (c) Mg-5CCA; and (d) Mg-7.5CCA composites.

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The grain size and grain morphology of Mg and Mg-CCA composites are shown in Table 1 and Figure 5. The grain morphology was found to be nearly equiaxed and the grain size decreased with increasing presence of reinforcement particles up to 5 wt %. Further increase in the amount of particles from 5 wt % to 7.5 wt % had no effect on grain size reduction and the average grain size remained the same in both the Mg-5CCA and Mg-7.5CCA composites. However, from the grain size distribution analysis, the grain size distribution was more homogeneous in Mg-7.5CCA’s composition (Figure 4d) when compared to Mg-5CCA’s composition (Figure 4c). While having the same average grain size in both compositions, the presence of small-sized grains was found to be more frequent in Mg-5CCA’s composition when compared to Mg-7.5CCA’s composition. In fact, the distribution graph resembles a right-skewed distribution in Mg-5CCA’s composition while a normal distribution pattern was observed in the Mg, Mg-2.5CCA, and Mg-7.5CCA composites.

Figure 4. Grain size distribution graphs of the: (a) Mg and (b) Mg-2.5CCA; (c) Mg-5CCA; and (d) Mg-7.5CCA composites. Table 1. Results of grain morphologyy and microhardness. Materials

Grain Size (μm)

Aspect Ratio

Microhardness (HV)

Mg Mg-2.5 wt % CCA Mg-5.0 wt % CCA Mg-7.5 wt % CCA

34 ± 4 14 ± 4 12 ± 5 12 ± 4

1.4 ± 0.3 1.4 ± 0.3 1.5 ± 0.3 1.5 ± 0.3

47 ± 2 56 ± 6 70 ± 6 80 ± 7

Since the reinforcement particle size used for the synthesis of Mg composites was at the micron length scale, particle-stimulated dynamic recrystallization during hot extrusion can be expected [21,22]. The more evenly distributed grain size (Figure 5) in the case of Mg-7.5CCA versus Mg-5HEA can be attributed to a more uniform distribution and a larger number of CCA particles in the former [23]. 51

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In extruded Mg composites, the recrystallized grains have common boundaries with the particles. With the more homogeneous distribution of particles throughout the matrix coupled with the increased presence of particles, it is possible that the migration of grain boundaries was prevented by the particles and hence the grain growth. This will provide more evenly distributed grains in the Mg-7.5CCA composite and hence while the largest grains observed in the Mg-5CCA composite measured in the range of 24–26 μm, it was only 20–22 μm in the case of the Mg-7.5CCA composite.

Figure 5. The grain size distribution pattern in the etched surfaces of the: (a) Mg; (b) Mg-2.5CCA; (c) Mg-5CCA; and (d) Mg-7.5CCA composites.

3.3. Mechanical Properties The result of microhardness measurements on Mg and Mg-CCA composites is shown in Table 1. It was observed that the indentation resistance of Mg was significantly increased with the presence of CCA particles. When compared to pure Mg, an increment in hardness of 1.2 times, 1.5 times, and 1.7 times was observed in the Mg-2.5CCA, Mg-5CCA, and Mg-7.5CCA composites, respectively. The localized matrix deformation was constrained by the presence of hard and strong ball-milled CCA particles. The increasing trend of hardness can be due to the addition of an increasing amount of hard CCA reinforcement particles in the Mg matrix. The compressive properties of pure Mg and Mg composites containing ball-milled CCA particles are listed in Table 2 and Figure 6. When compared to pure Mg, a significant improvement in compressive yield strength of 40%, 57%, and 63% in the Mg-2.5CCA, Mg-5CCA, and Mg-7.5CCA composites was achieved, respectively. In terms of ultimate compressive strength, 57%, 78%, and 79% increments over pure Mg were attained in the Mg-2.5CCA, Mg-5CCA, and Mg-7.5CCA composites, respectively. The results indicated an increasing trend of average compressive strengths with increased addition of ball-milled alloy particles from 2.5 to 7.5 wt % in the Mg matrix. Between the Mg-5CCA and Mg-7.5CCA composites, the strength increment was marginal and the strength level was the same 52

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if the standard deviation is taken into consideration. This can be explained based on the grain size measurement, which shows the occurrence of the same average grain size between the two composite compositions (Table 1). From the compressive failure strain results, the presence of CCA particles had no clear deteriorating effect on the ductility of pure Mg, which showed a similar failure strain to Mg and its composites (Table 2). However, the average failure strain in Mg-5CCA was found to be lower than that in Mg-2.5CCA and Mg-7.5CCA. From the grain size distribution result, a right-skewed distribution was observed with the influence of the small grain size in Mg-5CCA’s composition. For materials with coarse grains, a higher compressive strain can be expected due to continuous twinning within the coarse grains. For materials with smaller grains, the compressive strain can be decreased due to less continuity of twinning within the fine grains. This phenomenon was reported in detail in a related paper based on magnesium-based composites [16]. In addition, the ductility increment in this composite can be attributed to the resultant microstructural homogeneity in terms of grain size and particle distribution. The resultant lower failure strain in the Mg-5CCA composite can be accounted for with the influence of small grains from the grain distribution measurement (Figure 4c) and microstructural observation (Figure 5c). Table 2. Results of room temperature compressive properties.

Materials

0.2% Compressive Yield Strength (MPa)

Ultimate Compressive Strength (MPa)

Compressive Failure Strain (%)

Mg Mg-2.5 wt % CCA (1.6 vol %) Mg-5.0 wt % CCA (3.2 vol %) Mg-7.5 wt % CCA (4.9 vol %) Mg-6 vol % Ni50 Ti50 [17] Mg-5 vol % Ni60 Nb40 [18] AT81-5 vol % SiC [9] AZ91D-3 vol % TiC [10]

91 ± 8 127 ± 5 (40%) 143 ± 2 (57%) 148 ± 4 (63%) 89 ± 3 130 ± 11 127 ± 10 -

263 ± 16 414 ± 6 (57%) 469 ± 18 (78%) 472 ± 19 (79%) 368 ± 8 320 ± 11 301 ± 20 320 *

12 ± 2 15 ± 1 10 ± 2 15 ± 2 15.1 ± 1.5 18.4 ± 1.3 11.4 ± 0.5 17 *

* Values approximated from the Compression graph.

Figure 6. Compressive Stress-Strain curve of the Mg and Mg-CCA composites.

When compared to Mg composites containing ball-milled amorphous particles with a comparable amount of reinforcement addition, a significant improvement in compressive yield strength and ultimate compressive strength was observed in the Mg composite containing CCA particles

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while maintaining a similar compressive failure strain level (Table 2). In addition, significantly higher compressive strengths were achieved in the Mg-CCA composites when compared to the Mg-alloy-based composites containing micron size particles. This indicates the beneficial effect of CCA alloy particles for property enhancement in the composites under compressive loading. Furthermore, it shows the suitability and compatibility of this new type of CCA alloy particle as a reinforcement in an Mg matrix. Fracture surface studies were done on Mg and its composites and the representative fractographs are shown in Figure 7. From the fractographs, the appearance of smooth fracture surfaces can be seen in the Mg, Mg-2.5CCA, and Mg-7.5CCA composites. In case of the Mg-5CCA composite, the appearance of ragged and rough fracture features was observed. The observed fracture features conform with the resultant compressive failure strain values presented in Table 2, indicating the reduced ductility attained in the Mg-5CCA composite’s composition when compared to the Mg, Mg-2.5CCA, and Mg-7.5CCA composites (Figure 6).

Figure 7. Representative fractographs of: (a) Mg; and the (b) Mg-2.5CCA; (c) Mg-5CCA; and (d) Mg-7.5CCA composites.

4. Conclusions Based on the interrelation between the microstructural evolution and mechanical properties of the Mg-CCA composites developed in this work, conclusions are drawn as follows: 1. 2.

New Mg-CCA composites can be successfully developed using a powder metallurgy route incorporating microwave sintering and hot extrusion. The addition of ball-milled CCA reinforcement particles assisted in a significant refinement of the matrix grain size. The measurement on the grain size distribution showed a normal distribution in the Mg, Mg-2.5CCA, and Mg-7.5CCA composite compositions while a right-skewed distribution was observed in the Mg-5CCA composite.

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3.

4.

Hardness increased with an increasing amount of reinforcement addition in the Mg-CCA composites. The maximum microhardness of 80 HV was achieved in the Mg-7.5 wt % CCA composite. The compressive yield strength and ultimate compressive strength were significantly enhanced in the Mg-CCA composites while maintaining the same ductility levels as unreinforced Mg. The newly developed Mg-CCA composites showed higher strength under compressive loading when compared to Mg composites containing ball-milled amorphous particles and Mg-alloy-based composites containing micron-size particle reinforcement. The achievement of enhanced mechanical properties in Mg-CCA composites highlighted the effectiveness of using ball-milled CCA particles as a reinforcement in Mg.

Acknowledgments: The authors would like to acknowledge the Ministry of Education Academic Research Funding (WBS# R-265-000-586-114) for the financial support in carrying out this research work. Author Contributions: Manoj Gupta and Khin Sandar Tun proposed the original project and supervised the investigation. Yuming Zhang, Gururaj Parande, and Vyasaraj Manakari performed processing. Khin Sandar Tun, Yuming Zhang, Gururaj Parande, and Vyasaraj Manakari performed testing and characterization. Khin Sandar Tun performed data analysis and wrote the paper. Manoj Gupta contributed consultation, data analysis, and paper review. Conflicts of Interest: The authors declare no conflict of interest.

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Kulekci, M.K. Magnesium and its alloys applications in automotive industry. Int. J. Adv. Manuf. Technol. 2008, 39, 851–865. [CrossRef] Caton, P.D. Magnesium—An old material with new applications. Mater. Des. 1991, 12, 309–316. [CrossRef] International Magnesium Association. Available online: http://c.ymcdn.com/sites/intlmag.site-ym.com/ resource/resmgr/docs/automotive/MgShowcase15_Feb2011.pdf (accessed on 28 January 2018). Kainer, K.U.; Buch, F. The Current State of Technology and Potential for further Development of Magnesium Applications. In Magnesium Alloys and Technology; Wiley-VCH: Weinheim, Germany, 2003; pp. 1–22. Housh, S.; Mikucki, B.; Stevenson, A. Selection and Application of Magnesium and Magnesium Alloys. In ASM Handbook, 10th ed.; ASM International: Materials Park, OH, USA, 1990; Volume 2, pp. 455–479. Rohatgi, P.K. Metal Matrix Composites. Def. Sci. J. 1993, 43, 323–349. [CrossRef] Lloyd, D.J. Particle reinforced aluminum and magnesium matrix composites. Int. Mater. Rev. 1994, 39, 1–23. [CrossRef] Tjong, S.C. Novel nanoparticle reinforced metal matrix composites with enhanced mechanical properties. Adv. Eng. Mater. 2007, 9, 639–652. [CrossRef] Luo, D.; Pei, C.-H.; Rong, J.; Wang, H.-Y.; Li, Q.; Jiang, Q.-C. Microstructure and mechanical properties of SiC particles reinforced Mg–8Al–1Sn magnesium matrix composites fabricated by powder metallurgy. Powder Metall. 2015, 58, 349–353. [CrossRef] Cao, W.; Zhang, C.; Fan, T.; Zhang, D. In Situ Synthesis and Compressive Deformation Behaviors of TiC Reinforced Magnesium Matrix Composites. Mater. Trans. 2008, 49, 2686–2691. [CrossRef] Gupta, M.; Nai, S.M.L. Magnesium, Magnesium Alloys and Magnesium Composites, 1st ed.; John Wiley & Sons: Hoboken, NJ, USA, 2011; pp. 113–205. ISBN 978-0-47-049417-2. Perez, P.; Garces, G.; Adeva, P. Mechanical properties of a Mg–10 (vol %)Ti composite. Comp. Sci. Technol. 2004, 64, 145–151. [CrossRef] Tun, K.S.; Gupta, M.; Srivatsan, T.S. Investigating influence of hybrid (yttria + copper) nanoparticulate reinforcements on microstructural development and tensile response of magnesium. Mater. Sci. Technol. 2010, 26, 87–94. [CrossRef] Tun, K.S.; Gupta, M. Development of magnesium (yttria + nickel) hybrid nanocomposites using hybrid microwave sintering: Microstructure and tensile properties. J. Alloy. Compd. 2009, 487, 76–82. [CrossRef] Tun, K.S.; Gupta, M. Role of microstructure and texture on compressive strength improvement of Mg/(Y2 O3 + Cu) hybrid nanocomposites. J. Comp. Mater. 2010, 44, 3033–3050. [CrossRef]

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Tun, K.S.; Gupta, M. Compressive deformation behavior of Mg and Mg/(Y2 O3 + Ni) nanocomposite. Mater. Sci. Eng. A 2010, 527, 5550–5556. [CrossRef] Sankaranarayanan, S.; Hemanth Shankar, V.; Jayalakshmi, S.; Nguyen, Q.B.; Gupta, M. Development of high performance magnesium composites using Ni50 Ti50 metallic glass reinforcement and microwave sintering approach. J. Alloy. Compd. 2015, 627, 192–199. [CrossRef] Jayalakshmi, S.; Sahu, S.; Sankaranarayanan, S.; Gupta, S.; Gupta, M. Development of novel Mg–Ni60 Nb40 amorphous particle reinforced composites with enhanced hardness and compressive response. Mater. Des. 2014, 53, 849–855. [CrossRef] Jayalakshmi, S.; Gupta, M. Metallic Amorphous Alloy Reinforcements in Light Metal Matrices, 1st ed.; Springer: Cham, Switzerland, 2015; pp. 85–105. ISBN 978-3-319-15015-4. Tun, K.S.; Srivatsan, T.S.; Kumar, A.; Gupta, M. Synthesis of Light Weight High Entropy Alloys: Characterization of Microstructure and Mechanical Response. In Proceedings of the Twenty-Sixth International Conference on the Processing and Fabrication of the Advanced Materials (PFAM XXVI), Jeonju, Korea, 16–21 October 2017. Inem, B. Dynamic recrystallization in a thermomechanically processed metal matrix composite. Mater. Sci. Eng. A 1995, 197, 91–95. [CrossRef] Wang, X.J.; Wu, K.; Zhang, H.F.; Huang, W.H.; Chang, H.; Gan, W.M.; Zheng, M.Y.; Peng, D.L. Effect of hot extrusion on the microstructure of a particulate reinforced magnesium matrix composite. Mater. Sci. Eng. A 2007, 465, 78–84. [CrossRef] Chan, H.M.; Humphreys, F.J. The recrystallisation of aluminium-silicon alloys containing a bimodal particle distribution. Acta Metall. 1984, 32, 235–243. [CrossRef] © 2018 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

56

metals Article

A Meso-Mechanical Constitutive Model of Particle-Reinforced Titanium Matrix Composites at High Temperatures Weidong Song 1, *, Liansong Dai 1 , Lijun Xiao 1 , Cheng Wang 1 , Xiaonan Mao 2 and Huiping Tang 2 1 2

*

School of Mechatronical Engineering, Beijing Institute of Technology, Beijing 100081, China; [email protected] (L.D.); [email protected] (L.X.); [email protected] (C.W.) Northwest Institute for Non-ferrous Metal Research, Xi’an 710016, China; [email protected] (X.M.); [email protected] (H.T.) Correspondence: [email protected]; Tel.: +86-10-6891-4152

Academic Editor: Manoj Gupta Received: 31 October 2016; Accepted: 26 December 2016; Published: 7 January 2017

Abstract: The elastoplastic properties of TiC particle-reinforced titanium matrix composites (TiC/TMCs) at high temperatures were examined by quasi-static tensile experiments. The specimens were stretched at 300 ◦ C, 560 ◦ C, and 650 ◦ C, respectively at a strain rate of 0.001/s. scanning electron microscope (SEM) observation was carried out to reveal the microstructure of each specimen tested at different temperatures. The mechanical behavior of TiC/TMCs was analyzed by considering interfacial debonding afterwards. Based on Eshelby’s equivalent inclusion theory and Mori-Tanaka’s concept of average stress in the matrix, the stress or strain of the matrix, the particles, and the effective stiffness tensor of the composite were derived under prescribed traction boundary conditions at high temperatures. The plastic strains due to the thermal mismatch between the matrix and the reinforced particles were considered as eigenstrains. The interfacial debonding was calculated by the tensile strength of the particles and debonding probability was described by Weibull distribution. Finally, a meso-mechanical constitutive model was presented to explore the high-temperature elastoplastic properties of the spherical particle-reinforced titanium matrix composites by using a secant modulus method for the interfacial debonding. Keywords: titanium matrix composite; constitutive model; interfacial debonding; high temperature; elastoplastic properties

1. Introduction Titanium matrix composites (TMCs) become ideal materials for auto industry [1] and shipbuilding industry [2,3], with high specific strength, high specific modulus, and high temperature resistance. TMCs are mainly divided into two categories, continuously reinforced titanium matrix composites and particle-reinforced titanium matrix composites. Among them, particle-reinforced TMCs develop rapidly due to isotropic characteristics, high temperature properties, as well as low cost compared to the continuously-reinforced TMCs [4]. In order to achieve excellent properties, it is essential for reinforced particulates to have superior mechanical properties and also combine stably with the matrix materials [5]. Several ceramic particles were proposed as titanium reinforcements: SiC, B4 C, TiAl, TiB2 , TiN, TiC, and TiB [5–8]. Particularly, TiC was an excellent choice for its high modulus, strength, stiffness, hardness, and compatibility with titanium matrix [9,10]. According to the literature, TiC/Ti bulk nanocomposites have been significantly studied by Gu et al. [11,12], which systematically presented the influence of TiC on Ti matrix phase, densification, microstructure, and strengthening mechanisms. Metals 2017, 7, 15

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A series of theory about particulate-reinforced composites taking account of particle size effects, damage evolution and debonding damage was carried out by K. Tohgo [13–16]. The coupled effects of the temperature and strain rate were studied by Song et al. [17,18], and a modified Johnson-Cook model was proposed to predict the dynamic behavior of TiCp/Ti. Recently, the combination of macroscopic and mesoscopic methods have been widely used to investigate dynamic mechanical behaviors and constitutive model. Meso-mechanical damage theory considers the variation of stiffness/compliance tensor as one measure of the damage, so how to determine the effective elastic modulus of the damaged materials becomes a key problem. Main meso-mechanical theories are outlined as follows: Eshelby’s equivalent inclusion theory [19,20], self-Consistent theory [21,22], Mori-Tanaka’s theory [23], differential schemes [24], Hashin-Shtrikman Bounds [25,26], and so on. The elastoplastic behavior of the particle-reinforced composite with damage is widely explored by using the first-order stress moment, second-order stress moment, secant modulus method, and incremental method [27–29], the damage patterns include crack or hole in the matrix, interfacial debonding, particle fracture, and so on. In our previous work [30], a one-dimension dynamic constitutive model based on Eshelby’s equivalent inclusion theory and Mori-Tanaka theory was established, by adding micro-crack nucleation and growth model. A three-dimensional interfacial debonding model to predict the stress-strain responses of weakly bonded composites was proposed by Lissenden [31], which was based on a modified Needleman type cohesive zone model. Considering progressively weakened interface, an elasto-plastic multi-level damage model was developed to predict the effective elasto-plastic behavior of particle-reinforced metal matrix composites in the work of Lee and Pyo [32]. According to Xia and Wang [33], a micromechanical model based on the analysis of localized deformation bands was provided to predict the toughening of dual-phase composites. However, there is little literature on the elastoplastic behavior of particle-reinforced composite at high temperatures. Whether the models widely used at room temperature still effective at high temperatures in new materials still have not identified by research. In the current paper, the elastoplastic behavior of TiC particle-reinforced composite with interfacial debonding at high temperatures is discussed by means of Mori-Tanaka’s mean field theory in conjunction with Eshelby’s equivalent inclusion theory. A meso-mechanical constitutive model is proposed to predict the mechanical properties of the composite at high temperatures by considering the interfacial debonding. 2. Experimental Procedure 2.1. Materials The material of titanium matrix composite reinforced with 3% TiCp was provided by Northwest Institute for Nonferrous Metal Research, which was manufactured by the pre-treatment melt process. The composition of the titanium matrix alloy was Ti-6Al-2.5Sn-4Zr-0.5Mo-1Nb-0.45Si, which could be used at high temperature ranging from 600 ◦ C to 620 ◦ C with excellent strength and oxidation resistance maintained above 600 ◦ C. The reinforced particle dispersed homogeneously in the matrix which had an average diameter of about 5 μm [34] and no brittle phase existed. The interfacial reaction layers between the particle and the matrix were stable and the reaction zone width was below 3 μm, by which perfect ductility at room temperature and strength ratio above 650 ◦ C were demonstrated. 2.2. Specimen Preparation Specimens for quasi-static tensile tests were machined by linear cutting, which were in a shape of flat dumbbell with holes at both ends to be clamped. The thickness was 3 mm and the schematic of the specimens was presented in Figure 1.

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Metals 2017, 7, 15

Figure 1. Specimen for quasi-static tensile test (unit: mm).

2.3. Quasi-Static Tensile Tests The quasi-static tensile tests at elevated temperatures were carried on WDW-300 electronic universal testing machine (Jinan East Testing Machine Co., Ltd., Jinan, China). The temperature and measurement were controlled by GW-1200A controller and high-temperature furnace, respectively, during testing. The tests was conducted at deformation temperatures of 300 ◦ C, 560 ◦ C, and 650 ◦ C with the same strain rate of 10−3 s−1 . The heat should be preserved for 5–10 min to ensure a uniform temperature in the test piece after the specimens were heated to the experimental temperature. Each experimental condition was repeated at least three times, and the average was taken from two valid experimental data of good reproducibility to be the final result. SEM tests were performed by a BCPCAS4800 scanning electron microscope (JEOL Co., Ltd., Tokyo, Japan) to observe the fracture morphology of each specimen stretched at different temperatures. 2.4. Experimental Results 2.4.1. Microstructure The images in Figure 2 exhibit the fracture microstructure of TiC/TMCs composites samples tested at different temperatures. The results reveal that the failure of the composites is dominant by the interface debonding, particle cracking, and ductile fracture of the matrix. It can be seen that with the temperature rises up, the dimples of fracture surface tend to be more uniformly distributed. The sizes of dimples are getting to be larger and deeper when the experimental temperature increases from 300 ◦ C to 650 ◦ C. This demonstrates that the TiC/TMCs composites show better plasticity at elevated temperature.

(a)

(b)

(c)

Figure 2. SEM images of TiC/TMCs composites: (a) 300 ◦ C; (b) 560 ◦ C; and (c) 650 ◦ C.

2.4.2. Stress-Strain Relationship Figure 3 shows the results of the quasi-static stress-strain curves for the titanium matrix composite at different temperatures. According to the tests, it is obvious that the TiC/TMCs composites demonstrate temperature sensitivity. The flow stress decreases with increasing experimental temperature at the same strain rate. From the data at 300 ◦ C and 560 ◦ C, the typical strain hardening curve can be obtained, but the flow stress of the latter rises more slowly than the former. For the

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Metals 2017, 7, 15

stress-strain curve at 650 ◦ C, the flow stress dropped with increasing strain which indicates that the temperature softening effect is greater than strain hardening. As for the ductile properties, the composite exhibits better ductility with the increasing temperature, which means elongation is positively correlated with temperature.

Figure 3. Stress-strain curves for the TiC/TMCs composites under quasi-static tensile at different temperatures.

3. Meso-Mechanical Theory for Multiphase Composite 3.1. Average Stress of the Reinforcement and the Matrix Let us consider an n-phase composite, and we shall refer to the matrix as phase 0, the particle as phase 1, and damaged particle (void) as phase 2. Based on Mori-Tanaka theory, a uniform far-field stress σ is exerted on the composite and its boundary. Take the matrix as comparative material and then the average strain ε0 of the comparative material will satisfy Equation (1) under the same external force:   p σ = L 0 : ε0 − ε 0

(1)

p

where L0 and ε 0 are the stiffness tensor the plastic strain of the matrix, respectively. Due to the existence of the reinforcement, the average strain is actually different from ε0 and perturbation stress σ  and perturbation strain  ε are generated by the interaction between the reinforcement, so the average stress of the matrix is:   p σ(0) = σ + σ  = L 0 : ε0 +  ε − ε0

(2)

Since the elastic property and the coefficient of thermal expansion of the reinforcement is different from that of the matrix, the average stress of the reinforcement is expressed as:  + σ1 = σ(1) = σ + σ pt

  pt p L 1 : ε0 +  ε + ε1 − ε1 − α1∗   pt p = L0 : L0−1 : σ + ε + ε1 − Δε1 − α1∗ − ε1∗

60

(3)

Metals 2017, 7, 15

σ(2) = σ + σ  + σ2 = pt

pt

  pt p L 2 : ε0 +  ε + ε2 − ε2 − α2∗ = 0   pt p = L0 : L0−1 : σ + ε + ε2 − Δε2 − α2∗ − ε2∗

(4)

pt

where, σr and εr indicate the average stress difference and the average strain difference between the rth phase and the matrix, εr∗ is the eigenstrain of the rth phase, αr∗ is caused by the different coefficient of thermal expansion between the rth phase and the matrix: αrij∗ = (αr − α0 )ΔTδij = αr δij 

εr = S : εr∗ + Δεr + αr∗ pt

p



(5)

r = 1, 2

(6)

where αr is the thermal expansion coefficient of the rth phase. By using Equations (2), (3), (6), we have:   pt p σr = L0 : (S − I ) : εr∗ + Δεr + αr∗

r = 1, 2

(7)

where S is the Eshelby tensor and I is the four order unit tensor. The average stress of the composite meets the volume mixing ratio, namely: σ = f 0 σ(0) + f p σ(1) + f v σ(2)

(8)

where f 0 + f p + f v = 1. Substitute Equations (2) and (3) into Equation (8), we have:   pt pt σ  = − f p σ 1 + f v σ2 and:

(9)

 ε = ( I − S)[ f p (ε1∗ + Δε1 + α1∗ ) + f v (ε2∗ + Δε2 + α2∗ )] p

Let X =

ε1∗

p + Δε1

+ α1∗ , Y

=

ε2∗

p + Δε2

Y=

+ α2∗

p

(10)

and substitute Equation (10) into Equation (4), we get:

f p X + ( I − S)−1 L0−1 σ 1 − fv

(11)

Substituting Equation (10) into Equation (3), we have:

  −1

 f p X = L0 + ( L1 − L0 ) : S − 1−pf v (S − I ) : 1−1 f v ( L0 − L1 ) : L0−1 : σ + L1 : Δε1 + α1∗ 



   −1 p = 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : ( L0 − L1 ) : L0−1 : σ + (1 − f v ) L1 : Δε1 + α1∗

(12)

and: ε1∗ =



Y= ε2∗

    −1

  p : ( L0 − L1 ) : L0−1 : σ + 1 − f v − f p ( L0 − L1 ) : (S − I ) : Δε1 + α1∗ 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1

(13)

    −1 p : [( L1 − L0 ) : S + L0 ] : ( I − S)−1 : L0−1 : σ + f p L1 : Δε1 + α1∗ 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1

(14)







 −1

= 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : {[( L1 − L0 ) : S + L0 ] : ( I − S)      p − 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : Δε2 + α2∗

−1

:

L0−1

: σ + f p L1 :



p Δε1

+ α1∗



(15)

Substituting Equations (1), (10)–(12) into Equations (2) and (3), respectively, we get:   f σ(0) = σ + L0 : ( I − S) : f p X + f v Y = 1−1 f v σ + 1−pf v L0 : ( I − S)  

   −1

p : 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : ( L0 − L1 ) : L0−1 : σ + (1 − f v ) L1 : Δε1 + α1∗

61

(16)

Metals 2017, 7, 15     1− f − f σ(1) = σ + L0 : ( I − S) : − 1 − f p X + f v Y = 1−1 f v σ − 1−v f v p L0 : ( I − S)     −1

 p : 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : ( L0 − L1 ) : L0−1 : σ + (1 − f v ) L1 : Δε1 + α1∗

(17)

3.2. Effective Stiffness Tensor of the Composite The stress-strain relationship can be written as: ε = L −1 : σ

(18)

The average strain of the composite is the sum of the strain: ε = f 0 ε(0) + f p ε(1) + f v ε(2) Substitute ε(0) = ε0 +  ε − ε0 , ε ( 1 ) = ε0 +  ε + ε1 − ε1 − α1∗ and ε(2) = ε0 +  ε + ε2 − ε2 − α2∗ into Equations (18), (6), (10), we have: p

pt

p

pt

p

ε = L0−1 : σ + f p ε1∗ + f v ε2∗

(19)

Substituting Equations (13) and (15) into Equation (19), we get:

  ε = L0−1 : σ + f p ε1∗ + f v ε2∗ = L0−1 : σ + f p 1 − f v − f p [( L1 − L0 ) : S + L0 ]  −1   +f L : [( L − L ) : L0−1 : σ + 1 − f v − f p ( L0 − L1 ) : (S − I ) :  p 1  0 1   −1 p Δε1 + α1∗ + f v 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 :   p −1 −1 ∗ {[( L1 − L0 ) : S + L0 ] : ( I − S) : L0 : σ + f p L1 : Δε1 + α1 

    p − 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 : Δε2 + α2∗

(20)

The stiffness tensor of the composite at high temperatures can be derived from Equation (20):

   −1

: ( L0 − L1 ) : L0−1 L−1 = L0−1 + f p 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1   

  ∗ − 1 + 1 − f v − f p ( L 0 − L 1 ) : ( S − I ) : α1 : σ + f v 1 − f v − f p [( L1 − L0 ) : S    −1 p + L0 ] + f p L1 : {[( L1 − L0 ) : S + L0 ] : ( I − S)−1 : L0−1 + f p L1 : Δε1 + α1∗ : σ−1

   −1

= L0−1 + 1 − f v − f p [( L1 − L0 ) : S + L0 ] + f p L1 f p ( L0 − L1 ) : L0−1    + 1 − f v − f p ( L0 − L1 ) : (S − I ) : α1∗ : σ−1 + f v {[( L1 − L0 ) : S + L0 ] :  ( I − S)−1 : L0−1 + f p L1 : α1∗ : σ−1

(21)

4. Elastoplastic Analysis of the Composite 4.1. Constitutive Model of the Matrix The elastoplastic relationship of the matrix can be described by modified Ludwik equation:  n p (0) σ ( 0 ) = σ s + h ε0

(22)

(0)

where, σs is the yield stress of the matrix; h and n are material parameters determined by uniaxial tensile test. Under monotonic loading, the secant modulus E0s of the matrix is expressed as: E0s =

σ(0) p = + ε0

ε0e

1 (0)

σs E0 σ(0)

62

+

(23)

p

(0)

ε0

p n

σs + h(ε0 )

Metals 2017, 7, 15

p

where, ε0e and ε0 are the elastic strain and the plastic strain of the matrix, respectively. p Under three-dimensional stress, by replacing σ(0) and ε0 with Mises effective stress σ(0)∗ and p∗ effective strain ε0 , formula (22) can be rewritten as:  n p∗ (0) σ(0)∗ = σs + h ε0 where:

(24)

1

1

3 (0)  (0)  2 p ∗ 2 p (0) p (0) 2 σ(0)∗ = ( σij σij ) , ε0 = ( εij εij ) 2 3 (0) 

σij

(25)

is the stress deviator of the matrix. The secant bulk modulus and shear modulus of the matrix are expressed as: Es ks0 =  0 s  , 3 1 − 2v0

Es μ0s =  0 s  2 1 + v0

(26)

where v0s indicates the secant Poisson’s ratio. Due to plasticity incompressibility, the secant bulk modulus ks0 is equal to the elastic bulk modulus k0 , so, we have:   1 Es 1 v0s = − 0 (27) − v0 2 E0 2 In general, the elastoplastic behavior of the matrix under monotonic loading can be described by the secant Young’s modulus E0s and two elastic constants E0 and v0 . 4.2. Stress for the Reinforcement and the Matrix and the Secant Tensor of the Composite under Force Boundary Conditions When the matrix is in the elastoplastic stage, the modulus changes with the deformation, so the modulus of the matrix takes the secant value indicating by superscript S. According to Equations (16) and (17), the stress for the matrix and the reinforcement can be written as:

  f σ(0) = 1−1 f v σ + 1−pf v L0s : ( I − S) : 1 − f v − f p

     −1  s   p : L0 − L1 : L0s−1 : σ + (1 − f v ) L1 : Δε1 + α1∗ L1 − L0s : S + L0s + f p L1

(28)

: ( I − S ) : (1 − f v − f p ) 

 −1 p [( L1 − L0s ) : S + L0s ] + f p L1 : ( L0s − L1 ) : L0s−1 : σ + (1 − f v ) L1 : (Δε1 + α1∗ )

(29)

σ(1) =

1− f v − f p s 1 1− f v σ − 1− f v L 0

According to Equation (21), the secant tensor of the composite is given:  −1     −1  s  + 1 − f v − f p L1 − L0s : S + L0s + f p L1 f p L0 − L1 ( Ls )−1 = L0s   −1    

   : L0s + 1 − f v − f p L0s − L1 : (S − I ) : α1∗ + f v L1 − L0s : S + L0s   −1  : ( I − S)−1 : L0s + f p L1 : α1∗

(30)

L0s = (2ks0 , ks0 − μ0s , ks0 − μ0s , ks0 + μ0s , 2μ0s , 2μ0s )

(31)

L1 = (2k1 , k1 − μ1 , k1 − μ1 , k1 + μ1 , 2μ1 , 2μ1 )

(32)

L2 = (0, 0, 0, 0, 0, 0)

(33)

where:

 S=

I = (1, 0, 0, 1, 1, 1) αs

s

αs

s

αs

s

β β β 2 s − , − , + , βs , βs α, 3 3 2 3 2 3 2 63



(34) (35)

Metals 2017, 7, 15

1 + v0s , αs =  3 1 − v0s

βs =

  2 4 − 5v0s   15 1 − v0s

(36)

4.3. Interfacial Debonding Model In order to descript the propagation of the interfaces, Weibull statistical distribution is introduced to discuss the cumulative probability of the interfacial debonding. It is assumed that when the interfacial debonding happened between the matrix and the reinforcement (particle), the particle cannot bear any load and can be equivalent to a hole. By assuming that the tensile stress in the particle controls the debonding and the initial propagation strength along ii direction meets the Weibull statistical distribution Pii [35], we have: 

(1) Pii σ11



 

(1)

σ = 1 − exp − ii s

m i = 1, 2, 3

,

(37)

  (1) where, Pii σ11 is the ratio of the damaged particles to all particles, i.e., the debonding probability. s and m are scale parameter and shape parameter of the Weibull function. Thus, the volume fraction of the damaged particle on the composite is: 

(1) f 1 P11 σ11



⎧ ⎨

 

(1)

σ = f 1 1 − exp − 11 ⎩ s

m ⎫ ⎬ ⎭

(38)

The probability density of the damaged interface can be written as:  (1) m−1   (1) m   σ σ11 m (1) − 11 exp p11 σ11 = s s s

(39)

The relationship between the critical debonding strength of the interface σc and the two parameters s and m is given as [36]:    ∞ 1 (1) (1) σ11 pdσ11 = s·Γ 1 + (40) σc = m 0 When the critical debonding strength of the interface σc and the loading exerted on the particle are known, the volume fraction of the debonding particle can be obtained by Equation (37). 4.4. Elastoplastic Stress-Strain Relationship The damage constitutive relation can be expressed by Equation (30). When the matrix is in the plastic stage, L0s is not a constant and changes with the deformation process. At the same time, volume fraction f p and f v are also changing. So, in order to obtain the stress-strain relationship, L0s and f v of each stage should be calculated firstly. The numerical calculation is performed according to the following procedure: (1) calculate the effective elastic modulus L of the material and taking as the initial value; (2) for a given σ, determining σ(0) and σ(1) from Equations (28) and (29); (3) set σc and m, deriving f v from Equation (37) ( f p = f 1 − f v ) and then obtaining the effective stress of the (0)

matrix from σ(0) . If σ(0) is bigger than the elastic limit σs of the matrix, L0s should be calculated from Equations (23), (26), (27), (31); and (4) increasing σ and calculating Ls from the new L0s , then repeating the whole process. It is assumed that the elastic modulus of the TiC particle does not change with the changes in temperature and the elastic modulus of the matrix decreases with the increase of temperature. In order to obtain the elastic of the matrix at different temperature, we assume that the elastic modulus of the matrix decreases linearly with increasing temperature. Based the known elastic modulus of the matrix, the elastic modulus at different temperatures can be obtained. The elastic modulus and yield

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stress of the matrix at room temperature was experimentally determined. By assuming that there are linear relationships between elastic modulus and temperature, and yield stress with temperature, respectively, the elastic modulus and yield stress of the matrix at different temperatures was calculated, which was shown in Table 1. Table 1. The elastic modulus and yield stress of the matrix at different temperatures. Temperature

R.T.

300 ◦ C

600 ◦ C

650 ◦ C

700 ◦ C

E0 /GPa

113 1050

71 670

25 590

20 540

15 420

(0)

σs /MPa

5. Comparison of Numerical Predictions with Experimental Results When the uniaxial tensile stress σ11 is exerted on the particle-reinforced titanium matrix composite at high temperatures, the material parameters of the composite is listed as follows: ν0 = 0.35, (0)

E1 = 460 Gpa, ν1 = 0.188, f 1 = 0.03, h = 60 Mpa, n = 0.45, σc = 2.0σs , m = 5. The stress-strain curves of the composite at three different temperatures (T = 300 ◦ C, 560 ◦ C, 650 ◦ C) are shown in Figure 4.

Figure 4. Cont.

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Figure 4. Comparisons of stress-strain curves of TiC/TMCs composites with the theoretical results and test results at different temperatures: (a) 300 ◦ C; (b) 560 ◦ C; and (c) 650 ◦ C.

From Figure 4, it can be seen that the numerical predictions agree well with experimental results, which demonstrate the assumption that the elastic modulus decreases linearly with the increase of the temperature. The debonding model adopted in the current paper can be used to predict the elastoplastic behavior of the composite at elevated temperatures. There is little discrepancy of initial elastic modulus and yield stress, which may be caused by that the theoretical model did not consider the particle cracking and ductile fracture. Nevertheless, in the plastic section of the curve, the theoretical model prediction is not accurate enough, which needs further study. 6. Conclusions Based on Eshelby’s theory and Mori-Tanaka theory, the stress of the reinforcement and the matrix and the effective stiffness tensor of the composite under force boundary conditions are deduced. By using the assumption that the interfacial debonding is controlled by the tensile stress on the particle and the cumulative probability of the interfacial debonding is described by the Weibull function, a meso-mechanical constitutive model is proposed to investigate the elastoplastic properties of the particle-reinforced titanium matrix composite by using the secant modulus method. A good agreement between the numerical predictions and the experimental results is obtained, which demonstrate that the model and the method adopted in the current study is reliable and reasonable. When particle-reinforced titanium matrix composites were used at high temperatures, such as in aerospace and automobile industries, this model can be used to predict the mechanical properties so as to provide the theoretic basis for the design of structural parameters. Acknowledgments: The author acknowledges the financial support of the National Nature Science Foundation of China (11672043, 11521062, 11325209), the Opening Project of State Key Laboratory for Strength and Vibration of Mechanical Structures (SV2015-KF-10), the Opening Project of State Key Laboratory of Traction Power (TPL1701) and the Project of State Key Laboratory of Explosion Science and Technology (YBKT16-19, KFJJ15-02M). Author Contributions: Weidong Song conceived and designed the study; Liansong Dai and Lijun Xiao performed the experiments and analyzed the data; Weidong Song, Liansong Dai, Lijun Xiao and Cheng Wang made the theoretical analysis; Xiaonan Mao and Huiping Tang provided the specimens and essential parameters of the material; Weidong Song and Liansong Dai wrote the paper. All authors read and approved the manuscript. Conflicts of Interest: The authors declare no conflict of interest.

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Microstructures and Tensile Properties of Al–Cu Matrix Composites Reinforced with Nano-Sized SiCp Fabricated by Semisolid Stirring Process Feng Qiu 1,3 , Xiang Gao 1 , Jian Tang 1 , Yu-Yang Gao 1 , Shi-Li Shu 2 , Xue Han 3 , Qiang Li 1 and Qi-Chuan Jiang 1, * 1

2 3

*

Key Laboratory of Automobile Materials, Ministry of Education and Department of Materials Science and Engineering, Jilin University, Renmin Street NO. 5988, Changchun 130025, China; [email protected] (F.Q.); [email protected] (X.G.); [email protected] (J.T.); [email protected] (Y.-Y.G.); [email protected] (Q.L.) State Key Laboratory of Luminescence and Applications, Changchun Institute of Optics, Fine Mechanics and Physics, Chinese Academy of Sciences, Changchun 130012, China; [email protected] Department of Mechanical Engineering, Oakland University, Rochester, MI 48309, USA; [email protected] Correspondence: [email protected]; Tel./Fax: +86-431-85094699

Academic Editor: Manoj Gupta Received: 29 December 2016; Accepted: 3 February 2017; Published: 8 February 2017

Abstract: The nano-sized SiCp /Al–Cu composites were successfully fabricated by combining semisolid stirring with ball milling technology. Microstructures were examined by an olympus optical microscope (OM), field emission scanning electron microscope (FESEM) and transmission electron microscope (TEM). Tensile properties were studied at room temperature. The results show that the α-Al dendrites of the composites were strongly refined, especially in the composite with 3 wt. % nano-sized SiCp , of which the morphology of the α-Al changes from 200 μm dendritic crystal to 90 μm much finer equiaxial grain. The strength and ductility of the composites are improved synchronously with the addition of nano-sized SiCp particles. The as-cast 3 wt. % nano-sized SiCp /Al–Cu composite displays the best tensile properties, i.e., the yield strength, ultimate tensile strength (UTS) and fracture strain increase from 175 MPa, 310 MPa and 4.1% of the as-cast Al–Cu alloy to 220 MPa, 410 MPa and 6.3%, respectively. The significant improvement in the tensile properties of the composites is mainly due to the refinement of the α-Al dendrites, nano-sized SiCp strengthening, and good interface combination between the SiCp and Al–Cu alloys. Keywords: nano-sized SiCp ; aluminum matrix composites; mechanical properties; microstructures

1. Introduction In the past decades, particulate reinforced aluminum matrix composites (AMCs) have attracted much attention in the field of structural and functional materials [1–4]. SiCp reinforced AMCs have been a hot research issue in recent years because of their excellent properties such as low density, high tensile strength, high elastic modulus and wear resistance, etc. [5–7]. For example, SiCp reinforced AMCs are used for engine piston, and heat sink [8,9]. Compared with traditional micron-sized SiCp /Al composites, the higher tensile strength and good ductility of the nano-sized SiCp /Al composites entitle them to have more competitive ability for advanced structural applications such as in automotive and aerospace industries and the military [10]. In the past decades, several processing techniques have been developed for fabricating Al matrix composites reinforced with nano-sized particles such as high-energy milling, powder metallurgy, and nano-sintering, and liquid-state solidification processing (e.g., stir casting) [11–19]. In these techniques, the semisolid stirring process has some important

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advantages such as low cost, capability of producing products with complex shapes, and processing simplicity [14]. Al–Si and Al–Mg alloys are the usually used matrix phase in nano-sized SiCp /Al composites [10,16,20,21]. Xiong et al. [10] fabricated 14 vol. % nano-sized SiCp /Al–Mg composites. They reported that the ultimate tensile strength increased from 223 MPa to 286 MPa, while the ductility decreased from 4.8% to 3.9%. Hamedan et al. [16] produced 1.0 wt. % nano-sized SiCp /Al356 composite and reported that the ultimate tensile strength of the composite increased from 140 MPa to173 MPa, while the ductility decreased slightly from 6.1% to 5.38%. Compared to the Al–Si and Al–Mg alloys, the Al–Cu alloys can offer some good mechanical properties. For instance, a previous study by us has shown that the tensile strength and elongation increased by 26% and 50% respectively, for the modified Al–Cu alloys compared with unmodified alloy [22]. Moreover, we also found that the corrosion resistance of modified Al–Cu alloy had an improvement compared with the unmodified one [23]. However, to the best of our knowledge, so far, because there is no chemical affinity between the Cu element and SiCp , and the Cu element can also not improve the wettability between the aluminum matrix and SiCp [24], the Al–Cu alloys were rarely used as a matrix in the nano-sized SiCp /Al composites in stir casting. It is believed that if the nano-sized SiCp /Al–Cu composites with uniform distribution SiCp and clean interface between the SiCp and Al–Cu alloys could be successfully fabricated, the composites will exhibit excellent mechanical properties, which are very important for application in the automotive and aircraft industries. In this paper, the nano-sized SiCp /Al–Cu composites were fabricated by combining semisolid stirring with ball milling technology. Semisolid stirring can suppress the interfacial reaction due to the low stirring temperature [19]. The usage of precursor powders fabricated by the mix of the nano-sized SiCp and alloy powders using mechanical ball milling is of benefit to the dispersion of the nano-sized SiCp in the matrix due to the disruption of the agglomerate nano-sized SiCp clusters in advance. The microstructures and tensile properties of the synthesized composites were investigated, and the strengthening mechanism was discussed. We expect that such knowledge would provide guidance for the fabrication and application of the nano-sized SiCp /Al–Cu composites. 2. Experimental Procedure The Al–Cu alloy with a composition of (wt. %): 5.0 Cu, 0.8 Mn, 0.7 Fe, 0.5 Mg, 0.5 Si, 0.25 Zn, 0.15 Ti, 0.1 Cr and Al (balance) was used as the matrix. The nano-sized SiCp , with a purity of 99.9 wt. % and ~60 nm in diameter, were used as the reinforced particles. The morphology of the raw nano-sized SiCp particles is shown in Figure 1a. If the agglomerate nano-sized SiCp clusters are added into the melt directly, it is difficult for semisolid stirring to break the clustering and disperse the nano-sized particles uniformly. Figure 1b shows the Al–Cu alloy powders (99% pure) with average sizes of about 10 μm, their composition is the same as the Al–Cu alloy matrix. Figure 1c,d shows the precursor powders which are fabricated by the mix of the calculated nano-sized SiCp and Al–Cu alloy powders using mechanical ball milling with ZrO2 balls at the speed of 150 r/min for 50 h. Figure 1d is the a high magnification of the rectangular area in Figure 1c. It could be found that most of the nano-sized SiCp display a relatively uniform distribution in each individual composite particle surface. The ball to powder weight ratio was 8:1. During melting, Al–Cu alloy was melted at 933 K in air using an electricity resistant furnace and then cooled to 873 K at which point the matrix alloy was in semi-solid condition. The temperature range for the Al–Cu alloy used in this study to be in the semi-solid condition is 813 K–903 K. Then, the precursor powder was added into the molten metal after stirring the molten metal with a graphite stirrer at the speed of 500 r/min. After that, the melt was poured into a preheated steel die. After the casting process, the Al–Cu alloy and the composites were homogenised for 10 h at 758 K in order to avoid segregation. The materials were extruded to the batten shaped samples with the help of a 200-ton hydraulic press at 773 K with the extrusion ratio of 16. Before the tensile test, all the extruded samples underwent the T6 heat treatment (solutionized at 773 K for 2 h and aged at 433 K for 18 h).

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Figure 1. SEM images of morphologies of the raw (a) nano-sized SiC particles and (b) Al–Cu alloy powders; FESEM (field emission SEM) images of (c) nano-sized SiCp /Al–Cu composite powders after ball milling; (d) high magnification of the area marked in (c).

Microstructures of the composites were examined by an Optical Microscope (Axio Imager A2m, Zeiss, Oberkochen, Germany) equipped with image analysis software and a camera; a computer was used for the OM observation and the quantitative measurements of microstructural features. The size of Al dendrites in every composite was measured from forty images taken at two magnifications, such as 50×, and 100×. Five samples of every composite were used to obtain the standard deviations (the error bars) plotted in Figure 2. Microstructures of the composites and morphologies of the raw nano-sized SiC particles and Al–Cu alloy powders were observed by field emission SEM (FESEM, JSM6700F, Tokyo, Japan) and SEM (Evo18, Carl Zeiss, Oberkochen, Germany).

Figure 2. Grain sizes of α-Al in the cast Al–Cu alloy and nano-sized SiCp /Al–Cu composites with different SiCp contents.

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The extruded samples were machined into dog-bone shaped tensile samples with a gauge cross section of 5.0 mm × 2.5 mm and a gauge length of 30.0 mm. Tensile tests were conducted at room temperature by using a servo-hydraulic materials testing system (MTS, MTS 810, Minneapolis, MN, USA) at a constant strain rate of 3 × 10−4 s−1 . 3. Results and Discussion Figure 3 shows the as-cast microstructures of the Al–Cu alloy and nano-sized SiCp /Al–Cu composites with the nominal content of 1 wt. %, 3 wt. %, 5 wt. % SiCp . As shown in Figure 3a, the α-Al dendrites of the Al–Cu matrix alloy are coarse and their average size is about 200 μm. However, in the nano-sized SiCp /Al–Cu composites, the α-Al dendrites are significantly refined by the addition of nano-sized SiCp , as shown in Figure 3b–d. The refinement of the dendrite size is mainly due to some heterogeneous nucleation sties of the α-Al crystal provided by nano-sized particles during solidification, and the hindrance of the other added nano-sized SiCp to the growth of α-Al dendrites during the solidification process. Figure 2 shows the size of dendrite in the nano-sized SiCp /Al–Cu composites with different particle contents. In the 3 wt. % nano-sized SiCp /Al–Cu composite, the morphology of α-Al changes from coarse dendritic grain to equiaxial grain with finer sizes of about 90 μm, which increases the boundary concentration in the Al matrix. The increase in the boundary concentration could be helpful to improve the tensile strength of metals or alloys due to the grain boundary playing a role as a barrier to the transmission of the dislocations. The as-cast microstructure of the composite with 5 wt. % SiCp is similar to that with 3 wt. % SiCp , although the sizes of the α-Al dendrites were uneven sizes of 60–150 μm (Figure 3d). In the composite with 5 wt. % SiCp , the shape of Al dendrites became very non-uniform due to the agglomeration of nano-sized SiCp particles. The hindrance effect of the nano-sized SiCp on the α-Al dendrite growth is strong in the area of agglomeration of ceramic particles. On the contrary, the hindrance effect of the nano-sized SiCp on the α-Al dendrite growth is weakened in the area of less ceramic particles. Thus, the difference in the size of the α-Al dendrites was probably due to the nonuniform dispersion of SiCp when their contents reached 5 wt. %.

Figure 3. Cast microstructures of the nano-sized Al–Cu alloy and SiCp /Al–Cu composites with different SiCp contents; (a) Al–Cu alloy; (b) 1 wt. % SiCp ; (c) 3 wt. % SiCp ; (d) 5 wt. % SiCp .

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Figure 4 shows the engineering stress–strain curves of the cast Al–Cu alloy and nano-sized SiCp /Al–Cu composites, and Table 1 lists the detailed data of the tensile properties. As indicated, the strength and ductility of the composites (1 wt. % and 3 wt. %) are improved synchronously which is quite rare for the composites reinforced with ceramic particles, because in most reported works [11,14,20,21], the composites had a higher strength and lower ductility than the matrix alloy. The yield strength (σ0.2 ), ultimate tensile strength (UTS) and fracture strain (ε) of the nano-sized SiCp /Al–Cu composites firstly increase and then decrease with the increase in the content of SiCp . The 3 wt. % nano-sized SiCp /Al–Cu composite possesses the best tensile properties. The yield strength, UTS and fracture strain of the 3 wt. % SiCp /Al–Cu composite are 220 MPa, 410 MPa and 6.3%, which increase by 45 MPa (25.7%), 100 MPa (32.2%) and 2.2% (53.6%), respectively, compared to those of the as-cast Al–Cu alloy (175 MPa, 310 MPa and 4.1%).

Figure 4. Tensile stress–strain curves of the cast Al–Cu alloy and nano-sized SiCp /Al–Cu composites with different SiCp contents. Table 1. Tensile properties of the as-cast Al–Cu alloy and nano-sized SiCp /Al–Cu composites with different SiCp contents. SiCp (wt. %)

σ0.2 (MPa)

σb (MPa)

ε (%)

0 1 3 5

8 175+ −6 7 185+ −8 10 220+ −6 5 190+ −8

11 310+ −10 12 358+ −11 14 410+ −8 5 362+ −13

1.2 4.1+ −0.5 0.8 5.3+ −0.7 0.7 6.3+ −0.5 1.3 5.4+ −1.6

Figure 5a–d shows the FESEM images of the 3 wt. % and 5 wt. % nano–sized SiCp /Al–Cu composite, and TEM micrographs of the 3 wt. % SiCp /Al–Cu composite. As indicated in Figure 5a,b, more evenly distributed nano-sized SiCp particles in the 3 wt. % SiCp /Al–Cu composite are observed compared with the 5 wt. % SiCp /Al–Cu composite. However, in the 5 wt. % SiCp /Al–Cu composite as shown in Figure 5b, although there are some uniform distribution zones of nano-sized SiCp , the agglomeration of particles can still be easily found, as shown in Figure 5d. In other words, more and more particles aggregate to form the clusters with the increase in the content of SiCp , resulting in the quite uneven α-Al dendrites sizes and higher concentration of defects. As shown in Figure 5c, it is clearly seen that the nano-sized SiC particles dispersed inside the α-Al dendrite and the interface 73

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between the SiC particles and matrix are good and clean without any contaminations, indicating the superiority of this fabrication technology.

Figure 5. FESEM images of the (a) 3 wt. % and (b) 5 wt. % nano-sized SiCp /Al–Cu composite; (c) TEM micrographs of the 3 wt. % SiCp /Al–Cu composite; (d) the high magnification of area A in (b).

The significant improvement of the strength of the SiCp /Al–Cu composites is mainly due to the refinement of the α-Al dendrites and the hindrance of the nano-sized SiCp to the start and motion of dislocations in the matrix. Moreover, the significantly improved ductility of the nano-sized SiCp /Al–Cu composites with a simultaneously increased tensile strength is derived from three factors: (i)

Nano-sized reinforcement. Compared with micron-sized ceramic particles, the nano-sized ceramic particles used as reinforcement can not only possess higher tensile strength but also maintain good ductility, especially in the low contents [21]. Large reinforcement particles could give rise to cleavage in the particle due to the fact that they are acting as concentrators of stress, and lead to the formation of pits or cavities due to the loss of interphase cohesion. However, the smallest reinforcement particles usually do not initiate pits or cavities at the particle and bond well to the metal matrix [11]. (ii) Dendrite refinement. The refinement of the α-Al dendrites will result in the increase in matrix dendrite boundaries. The finer the dendrite is, the more tortuous the grain boundaries are. Therefore, the crack propagation becomes more and more difficult and thus the composites can endure the larger plastic deformation before fracture. (iii) Suppression of interfacial reaction. It is known that the reaction between molten Al and SiCp takes place easily in the temperature range from 675 ◦ C to 900 ◦ C, producing Al4 C3 which is a brittle and unstable phase [12]. The presence of Al4 C3 degrades the mechanical properties through crack propagation. In the present work, low stirring temperature (600 ◦ C) during the semisolid stirring process can suppress the interfacial reaction effectively, which will be helpful to restrict the formation of the Al4 C3 phase. The improved strength and cracking resistance of the interface bonding make the occurrence of the crack source cracking become more difficult.

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Nano-sized particle strengthening, microstructure refinement, and good interface between reinforcement and matrix with no brittle intermetallics can be responsible for the significant improvement in the mechanical properties of the nano-sized SiCp /Al–Cu composites. Hence, the nano-sized SiCp /Al–Cu composites showed high plasticity and strength. However, the properties of the composite with high content nano-sized SiCp particles could be weakened because this composite resulted in more agglomeration of the SiCp and higher concentration of defects. More severe agglomeration of SiCp could not lead to the matrix being completely wrapped up by the particles and thus result in the debonding of the interface. Moreover, micro-porosity and other defects around the SiCp clusters presented in the composites become the cracking source during the plastic deformation. The above analyses imply that the embrittlement of the composites resulting from micro-porosity and detects results in the decrease in strength and ductility of the 5 wt. % SiCp /Al–Cu. 4. Conclusions The nano-sized SiCp /Al–Cu composites with contents of 1 wt. %, 3 wt. %, 5 wt. % SiCp were successfully fabricated by combining semisolid stirring with ball milling technology. The α-Al dendrites are significantly refined due to the addition of nano-sized SiCp . The refinement of the dendrite size is mainly attributed to some nano-sized particles providing some heterogeneous nucleation sties of the α-Al crystal, and the hindrance of the other added nano-sized SiCp to the growth of α-Al dendrites during the solidification process. The strength and ductility of the composites are improved synchronously with the addition of nano-sized SiCp particles. The 3 wt. % nano-sized SiCp /Al–Cu composite displays the best comprehensive tensile properties, i.e., the yield strength, UTS and fracture strain increase from 175 MPa, 310 MPa and 4.1% of the as-cast Al–Cu alloy to 220 MPa, 410 MPa and 6.3%, respectively. Nano-sized particle strengthening, microstructure refinement, and a good interface between reinforcement and the matrix with no brittle intermetallics can be responsible for the significant improvement in the mechanical properties of the nano-sized SiCp /Al–Cu composites. Acknowledgments: The National Natural Science Foundation of China (NNSFC, No. 51571101), the “Thirteenth Five-year Plan” Science & Technology Research Foundation of Education Bureau of Jilin Province, China (Grant No. 2015-479), NNSFC (No. 51501176) and the Project 985-High Properties Materials of Jilin University. Author Contributions: Feng Qiu and Qi-Chuan Jiang conceived and designed the experiments; Feng Qiu, Xiang Gao, Jian Tang, Yu-Yang Gao and Qiang Li performed the experiments; Feng Qiu, Jian Tang, Shi-Li Shu, and Xue Han, analyzed the data; Feng Qiu wrote the paper. Conflicts of Interest: The authors declare no conflict of interest.

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Du, X.F.; Gao, T.; Liu, G.L.; Liu, X.F. In situ synthesizing SiC particles and its strengthening effect on an Al-Si-Cu-Ni-Mg piston alloy. J. Alloy. Compd. 2017, 695, 1–8. [CrossRef] Schöbel, M.; Altendorfer, W.; Degischer, H.P.; Vaucher, S.; Buslaps, T.; di Michiel, M.; Hofmann, M. Internal stresses and voids in SiC particle reinforced aluminum composites for heat sink applications. Compos. Sci. Technol. 2011, 71, 724–733. [CrossRef] Xiong, B.; Xu, Z.; Yan, Q.; Cai, C.; Zheng, Y.; Lu, B. Fabrication of SiC nanoparticulates reinforced Al matrix composites by combining pressureless infiltration with ball-milling and cold-pressing technology. J. Alloy. Compd. 2010, 497, L1–L4. [CrossRef] Kang, Y.C.; Chan, S.L.I. Tensile properties of nanometric Al2 O3 particulate-reinforced aluminum matrix composites. Mater. Chem. Phys. 2004, 85, 438–443. [CrossRef] Xiong, B.; Xu, Z.; Yan, Q.; Lu, B.; Cai, C. Effects of SiC volume fraction and aluminum particulate size on interfacial reactions in SiC nanoparticulate reinforced aluminum matrix composites. J. Alloy. Compd. 2011, 509, 1187–1191. [CrossRef] Hsu, C.J.; Chang, C.Y.; Kao, P.W.; Ho, N.J.; Chang, C.P. Al–Al3 Ti nanocomposites produced in situ by friction stir processing. Acta Mater. 2006, 54, 5241–5249. [CrossRef] Zhang, H.; Geng, L.; Guan, L.; Huang, L. Effects of SiC particle pretreatment and stirring parameters on the microstructure and mechanical properties of SiCp /Al–6.8Mg composites fabricated by semi-solid stirring technique. Mater. Sci. Eng. A 2010, 528, 513–518. [CrossRef] Mazahery, A.; Shabani, M.O. Characterization of cast A356 alloy reinforced with nano SiC composites. Trans. Nonferr. Met. Soc. China 2012, 22, 275–280. [CrossRef] Dehghan Hamedan, A.; Shahmiri, M. Production of A356–1 wt % SiC nanocomposite by the modified stir casting method. Mater. Sci. Eng. A 2012, 556, 921–926. [CrossRef] Nie, K.B.; Wang, X.J.; Wu, K.; Xu, L.; Zheng, M.Y.; Hu, X.S. Processing, microstructure and mechanical properties of magnesium matrix nanocomposites fabricated by semisolid stirring assisted ultrasonic vibration. J. Alloy. Compd. 2011, 509, 8664–8669. [CrossRef] Mazahery, A.; Abdizadeh, H.; Baharvandi, H.R. Development of high-performance A356/nano-Al2 O3 composites. Mater. Sci. Eng. A 2009, 518, 61–64. [CrossRef] Tahamtan, S.; Halvaee, A.; Emamy, M.; Zabihi, M.S. Fabrication of Al/A206–Al2 O3 nano/micro composite by combining ball milling and stir casting technology. Mater. Des. 2013, 49, 347–359. [CrossRef] Amirkhanlou, S.; Niroumand, B. Effects of reinforcement distribution on low and high temperature tensile properties of Al356/SiCp cast composites produced by a novel reinforcement dispersion technique. Mater. Sci. Eng. A 2011, 528, 7186–7195. [CrossRef] Yang, Y.; Lan, J.; Li, X. Study on bulk aluminum matrix nano-composite fabricated by ultrasonic dispersion of nano-sized SiC particles in molten aluminum alloy. Mater. Sci. Eng. A 2004, 380, 378–383. [CrossRef] Zhao, H.L.; Yao, D.M.; Qiu, F.; Xia, Y.M.; Jiang, Q.C. High strength and good ductility of casting Al-Cu alloy modified by Prx Oy and Lax Oy . J. Alloy. Compd. 2011, 509, L43–L46. [CrossRef] Xia, Y.M.; Bai, Z.H.; Qiu, F.; Jin, S.B.; Jiang, Q.C. Effects of multi-modification of rare earth oxides Prx Oy and Lax Oy on microstructure and tensile properties of casting Al-Cu alloy. Mater. Sci. Eng. A 2012, 558, 602–606. [CrossRef] Kobashi, M.; Choh, T. Effects of alloying elements on SiC dispersion in liquid aluminum. Mater. Trans. 1990, 12, 1101–1107. [CrossRef] © 2017 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).

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metals Article

Improved Compressive, Damping and Coefficient of Thermal Expansion Response of Mg–3Al–2.5La Alloy Using Y2O3 Nano Reinforcement Amit Kumar 1 , Khin Sandar Tun 1 , Amit Devendra Kohadkar 2 and Manoj Gupta 1, * 1 2

*

Department of Mechanical Engineering, National University of Singapore, 9 Engineering Drive 1, Singapore 117576, Singapore; [email protected] (A.K.); [email protected] (K.S.T.) Department of Mechanical engineering, Visvesvaraya National Institute of Technology, South Ambazari Road, Nagpur 440010, India; [email protected] Correspondence: [email protected]; Tel.: +65-6516-6358

Academic Editor: Daolun Chen Received: 6 March 2017; Accepted: 14 March 2017; Published: 21 March 2017

Abstract: In the present study, the effects of the addition of Y2 O3 nanoparticles on Mg–3Al–2.5La alloy were investigated. Materials were synthesized using a disintegrated melt deposition technique followed by hot extrusion. The samples were then characterized for microstructure, compression properties, damping properties, CTE (coefficient of thermal expansion) and fracture morphology. The grain size of Mg–3Al–2.5La was significantly reduced by the addition of the Y2 O3 nano-sized reinforcement (~3.6 μm, 43% of Mg–3Al–2.5La grain size). SEM and X-ray studies revealed that the size of uniformly distributed intermetallic phases, Al11 La3 , Al2 La, and Al2.12 La0.88 reduced by the addition of Y2 O3 to Mg–3Al–2.5La alloy. The coefficient of thermal expansion (CTE) was slightly improved by the addition of nanoparticles. The results of the damping measurement revealed that the damping capacity of the Mg–3Al–2.5La alloy increased due to the presence of Y2 O3 . The compression results showed that the addition of Y2 O3 to Mg–3Al–2.5La improved the compressive yield strength (from ~141 MPa to ~156 MPa) and the ultimate compressive strength (from ~456 MPa to ~520 MPa), which are superior than those of the Mg–3Al alloy (Compressive Yield Strength, CYS ~154 MPa and Ultimate Compressive Strength, UCS ~481 MPa). The results further revealed that there is no significant effect on the fracture strain value of Mg–3Al–2.5La due to the addition of Y2 O3 . Keywords: Mg–Al–RE alloy; magnesium alloy; damping; Al11 La3 phase; nanosize reinforcement; mechanical properties

1. Introduction Mg–Al-based alloys are considered important lightweight alloys due to their low density, high strength, and stiffness with good casting and processing ability. Although Mg–Al alloys exhibit a superior combination of mechanical properties, they are not suitable for application in automobile engine components due to their poor creep resistance [1,2]. It is well reported that poor creep properties in Mg–Al alloys are due to the formation of the β-eutectic phase (Mg17 Al12 ), which is unstable at high temperatures [3]. To improve the creep properties of Mg–Al, rare earth metals (RE) were used as alloying elements, as they can suppress the formation of the β-phase. In addition, RE also improved the grain refinement and strength while retaining the ductility, creep resistance, corrosion resistance and fatigue strength [4–7]. The addition of lanthanum (La) to Mg–4Al exhibited a good strengthening effect due to its precipitation hardening and grain refinement effects [3]. In our recent study on Mg–3Al–xLa (x = 1%, 2.5% and 4%), it was observed that the addition of La to Mg–3Al led to the consumption of most of the Al for the formation of Al11 La3 , Al2 La, and Al2.12 La0.88 intermetallic phases and

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suppressed the formation of the Mg17 Al12 phase [8]. Among all the compositions, the Mg–3Al–2.5La alloy exhibited the best tensile properties; Tensile Yield Strength, TYS ~160 MPa, Ultimate Tensile Strength, UTS ~249 MPa and fracture strain ~22%. However, the addition of La in Mg–3Al alloy caused a gradual decrease in the compressive strength and elongation [8]. On the other hand, nano-sized reinforcement (thermally stable ceramics such as Al2 O3 , ZrO2 , Y2 O3 ) used in magnesium-based nanocomposites has already shown potential improvement in the mechanical properties and ductility without any significant increase in the density [9–13]. Many types of advanced metal matrix nanocomposites are now easily available, and they exhibit functional properties. Recently, a few particle-reinforced, self-lubricating and self-healing metal matrix nanocomposites were synthesized using solidification techniques [14–16]. Hassan et al. [17] showed that the addition of nano-sized yttrium oxide (Y2 O3 ) particulates as a reinforcement in magnesium, synthesized by the disintegrated melt deposition (DMD) technique, enhanced the mechanical properties of the magnesium matrix. This work concluded that the addition of 1.9% Y2 O3 by weight exhibits the best mechanical properties compared to 0.6% and 3.1% Y2 O3 [17]. The present work addresses the further enhancement of the compression and damping response of Mg–3Al–2.5La alloy using Y2 O3 nano particulates as a reinforcement. Mg–3Al–2.5La alloy, containing 1.9% Y2 O3 by weight as reinforcement, is synthesized along with pure Mg, Mg–3Al and Mg–3Al–2.5La alloys, using the Disintegrated Melt Deposition (DMD) technique followed by hot extrusion. A detailed view of the effect of the Y2 O3 addition on the microstructure, Coefficient of Thermal Expansion (CTE), compression and damping properties of Mg–3Al–2.5La is provided. 2. Materials and Characterizations 2.1. Materials Magnesium turnings (99.9% purity) supplied by Acros Organics (Geel, Belgium) were used as the base material. Aluminium powder (99% purity) of size ~7–15 μm supplied by Alfa Aesar (Haverhill, MA, USA) and Mg–30%La master alloy supplied by Sunreiler Metal Co. Limited (Beijing, China) were used as alloying elements. Yttrium oxide (99.995% purity) of size 20–40 nm supplied by US Research Nanomaterials (Houston, TX, USA) was used as reinforcement in this study. 2.2. Processing Four different compositions, pure Mg, Mg–3%Al, Mg–3%Al–2.5%La and Mg–3%Al–2.5%La– 1.9%Y2 O3 by weight were synthesized using disintegrated melt deposition technique [18]. Pure Mg turnings, Al powder, Y2 O3 powder and Mg–30%La master alloy were placed in a multilayered sandwich fashion in a graphite crucible and superheated to 750 ◦ C under an argon gas atmosphere using electrical resistance furnace (Dakin Engineering Pte Ltd., Singapore). For uniform distribution of reinforcement particulates within the alloy matrix, the superheated slurry was then stirred at 450 rpm for 5 min using a stainless steel impeller (Starlight Tool Precision Engineering, Singapore) with twin blade (pitch 45◦ ). Stainless steel stirrer was used to avoid any iron contamination of the molten metal. After stirring, the molten melt was down poured through a nozzle of 10 mm diameter at the bottom of the crucible to the mould under the influence of gravity. Before entering the mold, the molten metal was disintegrated by two jets of argon gas, oriented normal to the melt stream. The flow of argon was maintained at 25 L/min [17]. An ingot of 40 mm diameter was then obtained. For synthesizing other compositions similar steps were followed. As cast ingot was later machined to 36 mm diameter and 45 mm length for the secondary processing. Secondary processing involved the soaking of ingot at 400 ◦ C for 1 h in a constant temperature furnace (Elite Thermal Systems Ltd., Market Harborough, Leicestershire, UK). Using a 150-ton hydraulic extrusion press, hot extrusion was carried out at 350 ◦ C die temperature with an extrusion ratio of 20.25:1 to obtain rods of 8 mm diameter. Extruded rods were further used to prepare samples for different characterization studies.

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2.3. Characterizations 2.3.1. Microstructural Characterization The microstructure was characterized using an optical microscope (Olympus Corporation, Shinjuku, Tokyo, Japan) on polished and etched samples (etchant: 4.2 gm picric acid, 10 mL acetic acid, 70 mL ethanol and 10 mL distilled water). The grain size was measured on the longitudinal section of samples, with the help of Scion image analysis software (beta 4.0.2, Frederick, MD, USA, 2000). To observe intermetallic phase formation and distribution, scanning electron microscopes JEOL JSM-6010 (JEOL Ltd., Tokyo, Japan) and Hitachi FESEM-S4300 (Hitachi, Ltd., Tokyo, Japan) equipped with energy dispersive spectrometric analysis (EDS) were used. X-Ray diffraction analysis was conducted using an automated Shimadzu LAB-XRD-6000 (Shimadzu Corporation, Kyoto, Japan) (Cu Kα:λ = 1.54056 Å) spectrometer with a scan speed of 2◦ /min. 2.3.2. Physical Characterization Density and Porosity: The density of extruded pure Mg, Mg–3%Al, Mg–3%Al–2.5%La and Mg–3%Al–2.5%La–1.9%Y2 O3 was measured using a gas pycnometer (Micromeritics Instrument Corp., Norcross, GR, USA). Each sample was run for five cycles to measure the density more accurately. Pure helium gas was purged with a pressure of 19.5 Psig for all the five cycles with a cycle fill pressure of 19.5 Psig. The difference between theoretical density (calculated by the rule of mixture) and experimentally measured density was quantified as the porosity level in the material. The Coefficient of thermal expansion: By using a thermo-mechanical analysis instrument LINSEIS TMA PT 1000LT (Linseis Thermal Analysis, Robbinsville, NJ, USA) the coefficient of thermal expansion (CTE) of pure Mg, Mg–3%Al, Mg–3%Al–2.5%La and Mg–3%Al–2.5%La–1.9%Y2 O3 was determined. The heating rate of 5 ◦ C/min was maintained with constant argon flow rate of 0.1 L per minute. The displacement of the test samples (each of 5 mm length and 8 mm diameter) was measured as a function of temperature (323 K to 673 K) using an alumina probe (Linseis Thermal Analysis, Robbinsville, NJ, USA). Damping: The vibrational damping capacity of the materials was measured using the resonance frequency damping analyzer (RFDA), (IMEC, Genk, Belgium). The vibration signal of each material (8 mm diameter, 60 mm length) was measured as a function of amplitude vs. time. 2.3.3. Mechanical Characterization Compression Properties: In accordance with ASTM E9-09, compressive properties of extruded pure Mg, Mg–3%Al, Mg–3%Al–2.5%La and Mg–3%Al–2.5%La–1.9%Y2 O3 samples were determined at ambient temperature, using a fully automated servo-hydraulic mechanical testing machine, MTS-810 (MTS systems corporation, Eden Prairie, MN, USA). The compression properties were measured at a strain rate of 8.334 × 10−5 s−1 . The specimens of 8 mm diameter, with length to diameter ratio of one were used. At least five different samples of each composition were tested to ensure repeatability of results. Fractured surfaces of all samples were analyzed using Hitachi S-4300 FESEM (Hitachi, Ltd., Tokyo, Japan). 3. Results and Discussion 3.1. Microstructural Characterization The microstructures of all the samples were initially characterized using SEM microscopy (JEOL Ltd., Tokyo, Japan) (Figure 1). Table 1 and Figure 2a show the grain size of different compositions after analysis. The results revealed that the addition of 2.5% La and 1.9% Y2 O3 to Mg–3Al reduced the average grain size by ~50%. It was observed that the addition of Al in Mg (Figure 1a,b) significantly reduced the grain size from ~22.6 μm to ~7.74 μm. It is frequently reported that during the solidification of Mg–Al alloys, fine grains are nucleated as the primary-Mg solid solution, along with

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the Mg17 Al12 eutectic mixture distributed along the grain boundaries [1,2,19,20]. Secondary processing or hot deformation during extrusion further breaks down the Mg17 Al12 network structure into fine precipitates, which results in grain refinement as observed in Mg–3Al (Figure 1b). Hot extrusion was performed at 350 ◦ C (which is >0.5 Tm of pure Mg), and therefore it resulted in recrystallization and the formation of nearly equiaxed grains. As the solubility of La in Mg is very limited (~0.78 wt %) [21,22] and La is a grain refiner, therefore the addition of 2.5% La to Mg–3Al alloy further reduced the average grain size from ~7.74 μm to ~6.26 μm. These results obtained in this study are in good agreement with other available reports claiming La as an excellent grain refiner in Mg [23,24]. The reinforcement of nano-sized thermally stable 1.9% Y2 O3 powder to Mg–3Al–2.5La further reduced the grain size as Y2 O3 nanoparticles can act as the nucleation sites during solidification and recrystallization besides pinning the grain boundaries in the later stages. Figure 2a represents the change in the grain size of pure Mg with the addition of 3Al, 2.5La, and 1.9Y2 O3 subsequently.

Figure 1. Scanning Electron Microscopic (SEM) micrographs of (a) pure Mg, (b) Mg–3Al, (c) Mg–3Al–2.5La, and (d) Mg–3Al–2.5La–1.9Y2 O3 alloys, illustrating the grain structure. Table 1. Results of average grain size, density, porosity and Coefficient of Thermal Expansion (CTE) measurements.

Material (wt %) Pure Mg Mg–3Al Mg–3Al–2.5La Mg–3Al–2.5La–1.9Y2 O3

Average Grain Size (μm) 22.6 ± 7.3 7.74 ± 1.5 6.26 ± 1.1 3.6 ± 0.5

Density and Porosity Measurements Theoretical Density (g/cc)

Experimental Density (g/cc)

Porosity (%)

CTE (×10−6 /K)

1.738 1.758 1.791 1.818

1.737 1.753 1.788 1.813

0.15 0.29 0.17 0.16

26.8 ± 3.9 26.1 ± 2.6 25.3 ± 2.7 25.0 ± 1.1

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(a)

        

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Figure 2. Change in (a) grain size and (b) coefficient of thermal expansion in pure Mg with the addition of 3Al, 2.5La and 1.9Y2 O3 subsequently.

Figure 3a–c are the SEM micrographs of Mg–3Al, Mg–3Al–2.5La, and Mg–3Al–2.5La–1.9Y2 O3 alloys. Figure 3a–b show the SEM micrographs of extruded Mg–3Al and Mg–3Al–2.5La alloys. In Mg–3Al alloy, the dispersed Mg17 Al12 phase is distributed inside the Mg matrix. A uniformly distributed bright white phase appeared in the Mg–3Al–2.5La alloy (Figure 3b) in rod-like (Al11 La3 ) and polygon-type (Al2 La, Al2.12 La0.88 ) shapes, which is consistent with earlier reports [3,21,25–29]. The brighter second phase in Figure 3c is broken into even finer shapes in the Mg–3Al–2.5La–1.9Y2 O3 alloy, especially rod-like shapes, illustrating the ability of Y2 O3 nanoparticles to refine the second phases. Similar findings were observed as a result of the addition of Al2 O3 in the AZ31 alloy [30]. The uniform distribution of the second phase is due to the hot extrusion, which broke down these scattered rod-like and polygon shapes into small pieces throughout the microstructure.

Figure 3. SEM micrographs of (a) Mg–3Al, (b) Mg–3Al–2.5La, and (c) Mg–3Al–2.5La–1.9Y2 O3 alloys; (d) compressive fractograph of Mg–3Al–2.5La–1.9Y2 O3 alloys.

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X-ray diffraction (XRD) studies conducted in the longitudinal direction of the samples are shown in Figure 4. These diffractograms did not reveal the presence of any La phase with Mg, which is consistent with other available reports on Mg–Al–La alloys [8,21,22,25–29,31,32]. However, they revealed the strong presence of Mg peaks together with the phase comprised of Al11 La3 , Al2 La and Al2.12 La0.88 , which are also observed in the SEM micrographs. The formation of Al11 La3 , Al2 La and Al2.12 La0.88 as intermetallic phases occurred due to the large difference in the electronegativity of Al and La when compared to Mg and Al [21,33,34]. It is well documented in the literature that dominating diffraction angles in extruded Mg rods corresponding to 2θ = 32◦ , 34◦ and 36◦ , respectively, represent the prismatic (1, 0, −1, 0) plane, the basal (0, 0, 2, 0) plane and the pyramidal (1, 0, −1, 0) plane of HCP Mg crystal [35]. From the intensity of these peaks at various diffraction angles, it is evident that the addition of La in Mg–3Al increased the I/Imax ratio for the basal plane but the pyramidal texture still dominated. The Mg–3Al–2.5La–1.9Y2 O3 alloy showed that the peak corresponding to the basal plane becomes dominant. This indicates that the presence of Y2 O3 clearly strengthens the basal texture in the Mg–3Al–2.5La alloy.

Figure 4. X-ray diffraction results of Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9Y2 O3 alloys.

3.2. Physical Characterization 3.2.1. Density and Porosity From Table 1, it is observed that near-dense Mg materials were synthesized utilizing the disintegrated melt deposition technique coupled with hot extrusion. The experimentally measured density values of the synthesized alloys and composite are closer to those of theoretically calculated density values. The increase in the density values of pure Mg was due to the addition of relatively high-density Al, La and Y2 O3 elements when compared to pure magnesium. The volumetric porosity results, which were calculated using theoretical and experimental density values, show that the addition of Y2 O3 did not affect the porosity of the base Mg–3Al–2.5La alloy. 82

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3.2.2. The Coefficient of Thermal Expansion Table 1 and Figure 2b show the results of the coefficient of thermal expansion (CTE) measurements within the 25–400 ◦ C temperature range. The results show that the CTE value of pure magnesium decreased with the alloying additions of Al, La and Y2 O3 . The gradual decrease in the CTE values of Mg, Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9Y2 O3 alloy was due to the presence of the alloying addition of Al, La and Y2 O3 which have lower CTE values (23.1 × 10−6 /K, 12.1 × 10−6 /K and 8.1 × 10−6 /K) as compared to pure Mg (26.8 × 10−6 /K) [36]. The results (see Figure 2b) suggest that the alloys and nanocomposites investigated in this study are more dimensionally stable with respect to temperature when compared to pure Mg. 3.2.3. Damping The damping characteristics of extruded pure Mg, Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9Y2 O3 alloys are presented in Table 2. The damping capacity of a material is defined as the ability to absorb vibration. The value of the damping capacity of a material depends on its properties such as density, microstructure, and elasticity [8,37]. The results show that the damping capacity of pure Mg decreased with the addition of Al in Mg–3Al, which was further enhanced by the addition of 2.5La in Mg–3Al–2.5La. The addition of Y2 O3 further improved the damping capacity of the Mg–3Al–2.5La. The damping loss rate represents how fast a material stops vibration. The results indicate that addition of 1.9% Y2 O3 decreased the damping loss rate compared to the addition of 2.5% La. The significant change in the damping properties of alloys can be due to the damping mechanisms related to texture reorientation, thermal mismatch, defects, porosity, dislocation and grain boundary. Table 2. Room-temperature compressive and damping properties of pure Mg, Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9Y2 O3 alloys. Material Pure Mg [8] Mg–3Al [8] Mg–3Al–2.5La [8] Mg–3Al–2.5La–1.9Y2 O3

0.2% CYS (MPa)

UCS (MPa)

Fracture Strain (%)

Damping Loss Rate

Damping Capacity

90 ± 6 154 ± 2 141 ± 4 156 ± 5

333 ± 4 481 ± 7 456 ± 3 520 ± 8

23 ± 0.74 24 ± 0.5 18 ± 1 18 ± 0.70

8.00 ± 1.000 6.16 ± 0.377 8.29 ± 0.827 7.60 ± 0.701

0.000456 0.000204 0.000265 0.000272

3.3. Mechanical Characterization Compression properties: Table 2 and Figure 5 show the room-temperature compression properties of extruded pure Mg, Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9Y2 O3 samples under compression loading. As evident from the results, the addition of 3Al in pure Mg enhanced the compressive yield strength (CYS), the ultimate compressive strength (UCS) and the fracture strain (FS) from ~90 MPa, ~333 MPa and ~23% to a level of ~154 MPa, ~481 MPa and ~24%. This increase in compressive strength was due to the hall-patch effect as there was a tremendous (~73%, ~22.6 μm to ~7.76 μm) reduction in grain size of the pure Mg. Another possible reason is the presence of fine Mg17 Al12 precipitates near the grain boundaries, which lead to precipitation hardening. The addition of 2.5La to Mg–3Al significantly reduced the CYS, UCS and failure strain values. In spite of the grain refinement (~7.74 μm to ~6.26 μm), the compression strength of Mg–3Al–2.5La decreased. This was due to the presence of intermetallic Al2 La and Al2.12 La0.88 with fine Al11 La3 phases, which are hard and exhibit sharp edges. Stress concentrates on these sharp edges and causes early crack initiation and subsequent crack propagation. The addition of nano-sized reinforcement particulates of Y2 O3 further refined the grain size of Mg–3Al–2.5La (~6.26 μm to ~3.6 μm) and fragmented the second phases, resulting in the best improvement of the CYS and UCS.

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Figure 5. Stress-strain curves of extruded pure Mg, Mg–3Al, Mg–3Al–2.5La and Mg–3Al–2.5La–1.9 Y2 O3 alloys under compression loading.

The presence of finer secondary phases assisted in restricting the motion of dislocations more effectively, leading to strength improvement in the case of the nanocomposite. Overall, the fracture strain remained unaffected (~18%) by the addition of the reinforcement when compared to the base alloy. Therefore, the addition of nano-sized Y2 O3 compensated for the decrease in the compression strength of Mg–3Al due to the addition of 2.5La, while the presence of La suppressed the formation of the Mg17 Al12 phase, which adversely affects the creep properties of Mg–3Al alloys. The compression fracture morphology of the Mg–3Al–2.5La–1.9Y2 O3 alloy is shown in Figure 3d. Compressive fractography studies (quasi-static) showed that the materials underwent the shear mode of deformation with the addition of the reinforcement. The ample split into two parts and the fracture surfaces of all samples were inclined at an angle of ~45◦ . The SEM fractograph of fractured surfaces revealed the presence of shear bands in the sample. Smooth fracture surfaces exhibited a ductile mode of fracture in the samples [38,39]. 4. Conclusions In this work, the effect of the addition of Y2 O3 on the microstructural and mechanical properties of Mg–3Al–2.5La alloy was primarily investigated. The following conclusions can be drawn: 1. 2.

3.

4.

With the addition of the Y2 O3 reinforcement, an even finer grain structure can be realized (~3.6 μm for Mg–3Al–2.5La–1.9 Y2 O3 alloy, 43% less than that of Mg–3Al–2.5La at ~6.26 μm). The microstructural characterization concluded that all intermetallic phases Al2 La and Al2.12 La0.88 and Al11 La3 were still present in dispersed form, but the sizes of these phases were refined by the addition of nanosize Y2 O3 in the Mg–3Al–2.5La alloy. The compressive results concluded that the addition of Y2 O3 to Mg–3Al–2.5La significantly improved the compressive yield strength and the ultimate compressive strength (CYS from ~141 MPa to ~156 MPa and UCS from ~456 MPa to ~520 MPa), which are even better than those of the Mg–3Al alloy (CYS, ~154 MPa and UCS, ~481 MPa). There was no adverse effect on the fracture strain value recorded for Mg–3Al–2.5La with the addition of Y2 O3 . The damping results concluded that the addition of nanosize Y2 O3 to Mg–3Al–2.5La improved the damping capacity. The addition of the Y2 O3 reinforcement also improved the CTE value of the Mg–3Al–2.5La alloy.

Acknowledgments: The authors would like to acknowledge the Ministry of Education Academic Research Funding (WBS# R-265-000-498-112) for the financial support in carrying out this research work. 84

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Author Contributions: Amit Kumar and Amit Devendra Kohadkar performed processing; Amit Devendra Kohadkar performed mechanical testing; Amit Kumar and Khin Sandar Tun performed microstructure studies and data analysis; Amit Kumar wrote the paper; Manoj Gupta contributed consultation, data analysis and paper review. Conflicts of Interest: The authors declare no conflict of interest.

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Characterization of In-Situ Cu–TiH2–C and Cu–Ti–C Nanocomposites Produced by Mechanical Milling and Spark Plasma Sintering Nguyen Thi Hoang Oanh 1, *, Nguyen Hoang Viet 1 , Ji-Soon Kim 2 and Alberto Moreira Jorge Junior 3,4,5,6,7 1 2 3 4 5 6 7

*

School of Materials Science and Engineering, Hanoi University of Science and Technology, No. 1 Dai Co Viet, Hanoi 100000, Vietnam; [email protected] School of Materials Science and Engineering, University of Ulsan, San-29, Mugeo-2 Dong, Nam-Gu, Ulsan 680-749, Korea; [email protected] Department of Materials Science and Engineering, Federal University of São Carlos, Via Washington Luiz, km 235, São Carlos, SP 13565-905, Brazil; [email protected] University of Grenoble Alpes, Science et Ingénierie des Matériaux et Procédés (SIMAP), F-38000 Grenoble, France Centre National de la Recherche Scientifique (CNRS), Science et Ingénierie des Matériaux et Procédés (SIMAP), F-38000 Grenoble, France University of Grenoble Alpes, Laboratoire d’Electrochimie et de Physico-chimie des Matériaux et des Interfaces (LEPMI), F-38000 Grenoble, France Centre National de la Recherche Scientifique (CNRS), Laboratoire d’Electrochimie et de Physico-chimie des Matériaux et des Interfaces (LEPMI), F-38000 Grenoble, France Correspondence: [email protected]; Tel.: +84-4-3868-0409

Academic Editor: Manoj Gupta Received: 5 February 2017; Accepted: 27 March 2017; Published: 29 March 2017

Abstract: This study focuses on the fabrication and microstructural investigation of Cu–TiH2–C and Cu–Ti–C nanocomposites with different volume fractions (10% and 20%) of TiC. Two mixtures of powders were ball milled for 10 h, consequently consolidated by spark plasma sintering (SPS) at 900 and 1000 ◦ C producing bulk materials with relative densities of 95–97%. The evolution process of TiC formation during sintering process was studied by using X-ray diffraction (XRD), scanning electron microscopy (SEM), and high resolution transmission electron microscopy (HRTEM). XRD patterns of composites present only Cu and TiC phases, no residual Ti phase can be detected. TEM images of composites with (10 vol % TiC) sintered at 900 ◦ C show TiC nanoparticles about 10–30 nm precipitated in copper matrix, most of Ti and C dissolved in the composite matrix. At the higher sintering temperature of 1000 ◦ C, more TiC precipitates from Cu–TiH2–C than those of Cu–Ti–C composite, particle size ranges from 10 to 20 nm. The hardness of both nanocomposites also increased with increasing sintering temperature. The highest hardness values of Cu–TiH2–C and Cu–Ti–C nanocomposites sintered at 1000 ◦ C are 314 and 306 HV, respectively. Keywords: spark plasma sintering; Cu–TiC; in-situ composites; mechanical milling

1. Introduction Metal matrix composites (MMCs) are advanced materials which combine ductility and toughness of metal and high strength and modulus of ceramic particles. The unique properties of MMCs are high specific strength, specific modulus, and good wear resistance compare to unreinforced metal [1]. In many type of MMCs, copper matrix composites (CMCs) have received a lot of interest because of super toughness and wear resistance which are used for structural application in wear industry [2].

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Generally, there are two routes to produce particulate-reinforced CMCs, which are ex-situ and in-situ. In the ex-situ method, ceramic particles such as TiB2, TiC, and oxide are introduced into the metal matrix via powder metallurgy or conventional casting methods [3,4]. However, the CMCs fabricated by these methods revealed a drawback because of poor interfacial bonding between reinforcement particles and copper matrix [5]. In order to improve the wettability of the Cu matrix and reinforcement phase, nano-ceramic particles were used [6]. Nevertheless, ceramic nanoparticles have tendency to segregate into clusters in milling process leading decrease strength of composite. The distribution of reinforced particles are non-uniform in the copper matrix, the mechanical and electrical features of the composite will be affected negatively [7]. On the contrary, ceramic particles synthesized by the in-situ method were dispersed more homogeneously in the copper matrix. The interfaces between reinforcement particles and matrix are clean, and very fine reinforcement particles are formed. Among CMCs, Cu–TiC system is attracted more attention due to their potential applications as electrical sliding contacts, resistance welding electrodes [8]. In the in-situ method, TiC nanoparticles were produced by the reaction between Ti and C during sintering process. In order to prevent grain growth of reinforcement and copper particles occur at high sintering temperature a fast sintering process need to be carried out. Spark plasma sintering (SPS) has some advantages such as rapid sintering, uniform sintering, low running cost, easy operation proves a suitable sintering technique for consolidation nano-structure, nanocomposite, and amorphous materials. In SPS, very high temperature over melting temperature may be attained in the contact area of powder particles which enhances interparticle bonding without considerable grain growth occurring [9–13]. The replacement of Ti powder in Cu–Ti–C composite by another powder such as TiH2 is considerable because of high price of Ti powder. In addition to, dehydrogenation of TiH2 occurs during sintering process is always accompanied by formation of high concentration of lattice defects and the highly activated Ti atoms. Released hydrogen from TiH2 will react with oxygen on the surface of TiH2 powders in the form of H2O which affect positively on the electrical conductivity of the composite [14]. The objectives of the present work are to explore the possibility of synthesizing Cu–TiC in-situ composites made from Cu–TiH2–C and Cu–Ti–C powder mixtures by mechanical milling and SPS. The effect of reinforcement content and sintering temperature on microstructure and hardness properties of composites was investigated. 2. Experimental Procedure The copper (with average particle size of 75 μm), titanium (average particle size of 45 μm), TiH2 (average particle size of 40 μm) and graphite (average particle size of 5 μm) powder (≥99% purity, from HIGH PURITY CHEMICALS Co., Ltd., Chiyoda, Japan) were used as starting materials. The powder mixtures of two composites Cu–TiH2–C and Cu–Ti–C with mixing ratio of 10 and 20 vol % TiC were mechanically milled in a high-energy planetary ball mill (P100-Korea). Milling was operated for 10 h at the rotational speed of 500 rpm and 0.5 wt % stearic acid was used as the milling process control agent. Balls and vials are made of stainless steel, the diameter of the balls was 5 mm and the powder-to-ball ratio was 1:10. The vial was evacuated and subsequently filled with argon up to 0.3 MPa. A 1.5 g amount of as-milled powder was loaded into a cylindrical graphite die with 10 mm-inner and was subjected to a pulsed current using a spark plasma sintering equipment, (SPS-515 apparatus Sumitomo Coal Mining, Tokyo, Japan). The chamber was pumped to low vacuum (1800 ◦ C) and plastic rupture [16], which could contribute substantially to the toughness of tungsten matrix. Riesch and Du et al. carried out extensive studies on small-size fabrication of tungsten fibre reinforced tungsten composites by chemical vapor infiltration or chemical vapor deposition [17–21]. Three-point bending and fibre pullout tests of the samples indicated that tungsten fibres could enhance the toughness of tungsten or tungsten alloy. However, on the economic aspect for low utilization rate of raw materials and harsh reaction conditions, chemical deposition method seems unfavorable. Composites of Wf /W have been prepared by powder metallurgical method, but the density of the materials needs to be improved [22,23] and other preparation methods need to be exploited. In our previous work [16,24], the spark plasma sintering (SPS) method was exploited to prepare tungsten-fibre reinforced tungsten composites (Wf /W), and compact composites were obtained because of the unique features of SPS (large pulsed DC current provides fast heating/cooling rate and short consolidation time, which helps obtain high density and fine grains). However, the non-uniform temperature distribution limits the large-scale production of Wf /W composites. Except for SPS, the hot pressing (HP) method is also used to consolidate composite materials, and is able to provide more uniform heat distribution and enables preparation of large-size samples. In addition, cold rolling is a promising method to enhance material properties due to the porosity reduction and texture evolution in the rolling process [25], thus the cold rolling after HP (HPCR) may further improve the performance of the composites. In this work, different preparation methods based on SPS, HP and HPCR were exploited to prepare tungsten-fibre-net reinforced Wf /W samples. The influences of preparation method on the microstructure, Vickers hardness and tensile properties of the Wf /W composites were comparatively investigated. 2. Experimental Section 2.1. Preparation and Characterization Commercial K-doped tungsten fibres with diameter of 150 μm were purchased from Honglu Molybdenum Company, Xiamen, China. The chemical composition of the fibres is given in our previous work [16,24]. The theoretical density of the W fibres can be considered as 19.3 g/cm3 , as same as the value of pure tungsten. Long tungsten fibres were woven into a net-like shape using a braider, as presented in Figure 1. The fibres along the YO direction act as supporting effect of the nets, while the fibres along the XO direction (which is also the tensile and rolling direction) will play the major role in improving the toughness. The diameter of the nets is about 20 mm for SPS as shown in Figure 1a and about 60 mm for HP and HPCR as shown in Figure 1b. The enlarged picture at the part I in Figure 1b is shown in Figure 1I for clarity. In this work, each sample contained four tungsten-fibre-net layers according to the optimization of fibre contents in our earlier work [24].

Figure 1. Photograph of a tungsten fibre net for SPSed samples (a), HPed and HPCRed samples (b), and the enlarged picture at the part I in b (I).

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Pure tungsten powders (purity > 99.9%, average particle size 600 nm) were ball-milled for 4 h in a planetary ball mill in argon atmosphere to increase the activity with a powder-to-ball weight ratio of 1:8 and rotation speed of 240 rpm. During ball milling, tungsten carbide balls and mortars were used to minimize the possible contamination by impurities. Tungsten fibre nets were buried in the tungsten powders. To keep a roughly equal distance between the net layers (about 750 μm), equal amount of W powder was put into the space between the neighboring net layers. The fibre mass fraction in the composite was calculated just by weighting the fibres and matrix powders respectively before mixing. The samples with the diameter of 20 mm were sintered by spark plasma sintering (SPS, FCT systeme GmbH, Rauenstein, Germany) according to the procedure as described in [26], which is shown in Figure 2a again for clarity, where the temperature was 1800 ◦ C and the pressure was 47 MPa. The samples with the diameter of 60 mm were sintered in a hot-pressing vacuum furnace (HP, Shanghai Chen Xin Electric Furnace Co., Ltd., CXZT-60-23Y, Shanghai, China) with a heating or cooling rate of about 10 ◦ C/min. The detailed temperature and pressure profile of the HP sintering program were illustrated in Figure 2b: (i) heated from room temperature to 700 ◦ C and held for 60 min; (ii) heated to 1300 ◦ C and held for 180 min; (iii) heated to the sintering temperature 1800 ◦ C and held for 90 min; and (iv) cooled down to room temperature (RT). At the same time, the pressure in the graphite die was increased linearly to 56 MPa with a changing rate of about 0.28 MPa/min, held for 300 min and then decreased to zero. After that, The HPed samples were preheated in a high temperature furnace (KSL-1100X, Hefei Kejing Materials Technology Co., Ltd., Hefei, China) at 1100 ◦ C for 5 min, and then rolled by a rolling machine (Wuxi Guancheng Machinery Co., Ltd., Wuxi, China) with the roller radius of 100 mm, the roller length of 200 mm and the rotating speed of 500 r/min [27]. The samples were also preheated between the successive rolling. The HPed plates were rolled three times with a total thickness reduction of 38%. Hereafter, the cold rolled HPed samples were abbreviated as HPCRed ones.

Figure 2. Temperature and pressure profile of SPS and HP procedure: (a) SPS and (b) HP.

2.2. Testing Procedure The density of samples was determined by Archimedes principle. The sintered samples and the inlaid original wire were polished and then chemically etched with a 10% aqueous solution of K3 Fe(CN)6 and NaOH. The distribution of the tungsten fibres in the composites was investigated by optical microscopy (ZEIZZ-Axio Scope.Al, Carl Zeiss AG, Oberkochen, Germany). The Vickers micro-hardness of the sample was tested by the Vickers Indenter (HV-1000 A, Laizhou ITC Test Instrument Co., Ltd, Laizhou, China) at room temperature with a load of 200 g and a dwell time of 10 s. The hardness of matrix was tested along the direction either perpendicular or parallel to the pressing direction, either close to or away from the fibre. Each sample was indented for 8 times in different locations and the average value was adopted.

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In order to obtain information about the grain size, aspect ratio of grains, angles of grain boundary and anisotropy or isotropy of fibre and matrix in samples, the Electron Back Scattering Diffraction (EBSD) measurements were carried out. The samples for EBSD were mechanically polished and then electro polished by NaOH (2%) aqueous solution. The EBSD data was collected and analyzed using the software of HKL Tango (Version-2010, Oxford Instruments, Oxford, UK) with the resolution rate of 80–90% and a step from 0.25 to 0.4 μm depending on the grain size, and the raw data was presented. Microstructure of the samples and the fracture surfaces of the tensile tested samples were characterized with a field-emission scanning electron microscope (FESEM Sirion 200, FEI, Hillsboro, OR, USA). The dog-bone-shaped tensile samples were cut along the XO direction with a working length of 5 mm as in [9] and a cross-section of 1.4 × 1.5, 1.4 × 2.2, 1.4 × 1.8 mm2 for the SPS, HP, and HPCR samples, respectively. Tensile experiments were carried out at various temperatures using an Instron-5967 machine (Instron Corporation, Boston, MA, USA) with a constant displacement rate of 0.06 mm/min. In this work, the loading direction is parallel to the XO direction due to high mechanical performance [28], the number of tested samples was three to five. 3. Results and Discussion 3.1. Density and Vickers Hardness The mass fraction of W fibres, relative density and Vickers hardness of the Wf /W composites are listed in Table 1. The estimated mass fraction of fibres in SPSed, HPed and HPCRed samples is 17.4%, 10.5% and 10.5%, respectively. As listed in Table 1, the relative densities of the SPSed (97.5%) and HPed composites (95.1%) are a little lower than that of the HPCRed samples (99.8%), which is close to 100% of pure dense tungsten. The overall average hardness of fibres in composites is approximately 537 ± 7 HV0.2 for all three kinds of samples, which is about 100 HV0.2 smaller than that of the original tungsten fibre (623.6 ± 11 HV0.2 ) due to the grain growth of fibre in the sintering process [29], and more details will be analyzed in the following sections. Indicated by the lower density, the hardness of the matrix in the HPed samples (331.3 ± 8 HV0.2 ) is lower than that of the SPSed (431.3 ± 8 HV0.2 ) and HPCRed samples (488.2 ± 810 HV0.2 ). Table 1. Density, Mass fraction of fibres and Vickers micro-hardness of the fibre and matrix in the SPSed, HPed and HPCRed samples.

Different Samples

Mass Fraction of Fibre

Relative Density (%)

Original fibre SPSed HPed HPCRed

100% 17.4 ± 0.1% 10.5 ± 0.1% 10.5 ± 0.1%

100% 97.5 ± 0.3% 95.1 ± 0.2% 99.8 ± 0.1%

Vickers Hardness/HV0.2 Fibre

Matrix

623.6 ± 11 537.4 ± 9 538.5 ± 7 536.4 ± 5

431.3 ± 8 331.9 ± 6 476.4 ± 2

3.2. Distribution of Fibres To investigate the distribution of the tungsten-fibre-nets in the SPSed, HPed and HPCRed samples, cross-sectional optical micrographs of the samples with a low magnification are shown in Figure 3. Due to different grain size and grain orientation, the fibres and matrix show different corrosion extent by the etching solution. All the darker areas indicate the tungsten fibres, where the wave strips are the fibres in YO orientation while the round circles represent fibres in XO orientation. From Figure 3a it can be seen that the fibres distribute neatly in the net layer of the SPSed samples, and the net layers are separated one by one with a nearly equal distance. In the HPed samples (Figure 3b), the fibre layers are also evenly and regularly aligned. After cold rolling, the HPCRed samples are much thinner so that a slight variation in the fibre distribution and orientations can be obtained from the tomographic visualization as shown in Figure 3c. Before the rolling phase of preparation, the fibres have a circular shape with an interlayer spacing of about 560 μm. After rolling, the fibres present elliptical shape with 135

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an interlayer spacing of about 350 μm. The 38% reduction in spacing is consistent with the thickness reduction of the whole sample.

Figure 3. Optical micrographs of the Wf /W composites: (a) SPSed, (b) HPed, and (c) HPCRed.

The SEM images are presented in Figure 4 for (a) SPSed, (b) HPed and (c) HPCRed samples fractured at RT. It can be seen that there is no debonding seen in any of the samples between the tungsten fibre and the matrix. In addition, the cross section of fibre in the SPSed and HPed samples is almost circular, while that in the HPCRed samples is elliptical with an aspect ratio of about 5:2 owing to the rolling.

Figure 4. The Scanning Electron Microscope (SEM) images of the Wf /W composites: (a) SPSed, (b) HPed, and (c) HPCRed.

3.3. Microstructure of the Original and Sintered Fibres Figure 5a shows the EBSD results of the original and sintered W fibres, where the black lines mean the high-angle grain boundaries (θ > 10◦ ) while the gray lines mean the low-angle grain boundaries (θ < 10◦ ). Figure 5(a1) shows the EBSD results of the original W fibres. The ZO direction is perpendicular to the XO and YO directions. The texture of the original tungsten fibre is as shown in Figure 5(a2). Figure 5(a1) also indicates that the grains of the original tungsten fibre are tens of microns long in the XO direction and several microns wide in the YO and ZO directions, corresponding to a length/width aspect ratio of about 5.2:1 (as shown in Figure 5(a3)). It is interesting to note that the elongated mother-grains (θ > 10◦ ) are composed of fine equiaxed sub-grains (θ < 10◦ ) with the grain size less than 1 μm. In addition, the ratio of small angle grain boundaries in the original fibres is as high as 60%.

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Figure 5. The Electron Backscattered Diffraction (EBSD) images of a fibre in original W fibres-a, SPSed-b, HPed-c, HPCRed-d samples: (1) grain boundary and Euler angle map, (2) the inverse pole figures, (3) length-to-diameter ratio, and (4) grain boundary misorientation map.

It is well known that the mechanical properties are determined by the microstructure of materials, so it is necessary to investigate the microstructure of a fibre and matrix in the sintered samples. The microstructure of fibres in the SPSed, HPed and HPCRed samples is clearly demonstrated by the EBSD results as shown in Figure 5b–d, respectively. According to the deformation texture analysis, the grain width at the edge of the fibre in SPSed samples is a little larger than that in the central, as shown

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in Figure 5(b1). However, the texture of fibre in SPSed samples is , which is consistent with the original fibre, as shown in Figure 5(b2). On the other hand, the average aspect ratio of grains with misorientation angles θ > 10◦ is about 3.7:1 (Figure 5(b3)). From the grain boundary misorientation map shown in Figure 5(b4), it can be noted that the percentage of grains with misorientation angle 10◦ is about 3.2:1, as shown in Figure 5(c3), which is a little smaller than that in the SPSed samples. The percentage of grains with misorientation angles θ < 10◦ decreases to 26.8%, indicating that partial recrystallization occurred in the HP sintering process more severely than in the SPS process. For the HPCRed samples, however, the texture density of fibres in the XO direction decreases, and the orientation disperses along in both the YO and ZO directions as shown in Figure 5(d1,d2). According to Figure 5(d3), the grains become longer and the average aspect ratio is about 5.8:1. Furthermore, the proportion of grains with θ < 10◦ is increased to 31.6% (Figure 5(d4)). All the variation of the grains in the fibres could be attributed to the rolling process. As well known, high temperature annealing can result in reduction of dislocation density and elimination of some low-angle grain boundaries [30]. Therefore, during the high temperature sintering, stress in the grain of fibre is gradually released and the number of small angle grain boundaries decreases due to grain growth and coalesce. As a result, the percentage of grains with misorientation angle 10◦ and the percentage of grains with a small misorientation angle (θ < 10◦ ) decreases to 3.6 μm and 4.2%, respectively. For the HPed samples, the EBSD results of W matrix near and far away from tungsten fibre were shown in Figure 7. Comparing Figure 7(a1) with Figure 7(b1), there is no obvious texture, but the grain size is much different. The average size of grains with large misorientation angle (θ > 10◦ ) is 8.1 μm near the tungsten fibre and 11.4 μm far away from the fibre. In addition, the percentage of grains with small misorientation angle (θ < 10◦ ) near and far away from tungsten fibre in the HPed samples is 7.8% and 12.0%, respectively.

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Figure 6. EBSD images of tungsten matrix near the fibre-a, far away from the fibre-b in SPSed samples: (1) grain boundary and IPF map (inset), (2) grain size, and (3) grain boundary misorientation map.

Figure 7. EBSD images of tungsten matrix near the fibre-a, far away from the fibre-b in HPed samples: (1) grain boundary and IPF map (inset), (2) grain size, and (3) grain boundary misorientation map.

Figure 8 shows the EBSD results of tungsten matrix near and far away from tungsten fibre in the HPCRed samples, which is indicated by white (high-angle grain boundaries) and black lines (low-angle grain boundaries). Intuitively, tungsten grains in the matrix are elongated along the rolling direction in the HPCRed samples and the average grain length/width ratio is about 2.6:1 near the fibre (see Figure 8(a2)) and 2.1:1 away from the fibre (see Figure 8(b2)). There is also a large number of small angle grain boundaries and the proportion of grains with small misorientation angles (θ < 10◦ ) is about 72.7% near the fibre (see Figure 8(a3)) and 78.3% away from the fibre (see Figure 8(b3)). To sum up, the EBSD results of all samples were shown in Table 2. It can be seen that the W fibres embedded in the W matrix have great influence on the microstructure of the tungsten matrix during different sample preparation. In the SPS sintering process, because of the non-uniform current distribution and the higher density of fibre than matrix, the temperature is higher in and near the fibre, which can lead to grain growth in the fibre as shown in Figure 6(a1), and result in the lower proportion of small angle grain boundaries (θ < 10◦ ) and the larger grain size near the fibre. However, the exposure time to high temperature is not long enough to make the grains change from the strip-like to the equiaxed in tungsten fibres. In the HP sintering process, although the temperature is more uniform, the long exposure time to high temperature can result in a lower percentage of small angle grain boundaries in the fibre and the larger grain size in the matrix owing to the recrystallization.

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In the HPCRed samples the plenty of small angle boundaries in the matrix both near and far away from the fibre and the small grain size of W matrix can be attributed to the severe plastic deformation and dynamic recrystallization during rolling, as in the cold rolled pure W samples [31].

Figure 8. EBSD images of tungsten matrix near the fibre-a, far away from the fibre-b in HPCRed samples: (1) grain boundary and IPF map, (2) length-diameter ratio, and (3) grain boundary misorientation map. Table 2. Small angle grain boundaries (θ < 10◦ ) and grain size of the original fibre, the fibre and matrix in SPSed, HPed and HPCRed samples. The Share of Small Angle Grain Boundaries θ < 10◦ /%

SPSed HPed HPCRed

Average Length-Diameter Ratio

60.1 ± 0.2

Fibre

5.2 ± 0.1

Fibre

Near the Fibre

Away from the Fibre

Fibre

Near the Fibre

Away from the Fibre

31.5 ± 0.1 26.8 ± 0.2 31.6 ± 0.3

6.8 ± 0.2 7.8 ± 0.1 72.7 ± 0.3

4.2 ± 0.1 12.0 ± 0.2 78.3 ± 0.2

3.7 ± 0.1 3.2 ± 0.1 5.8 ± 0.2

4.3 ± 0.3 μm (grain size) 8.1 ± 0.2 μm (grain size) 2.6 ± 0.1

3.6 ± 0.2 μm (grain size) 11.4 ± 0.2 μm (grain size) 2.1 ± 0.2

3.5. Tensile Properties and Fracture Microstructure Tensile tests were performed at different temperatures for the SPSed, HPed and HPCRed samples, and the curves of engineering stress versus strain were shown in Figure 9. From such curves the average tensile strength (TS) and total elongation at break (TE) at different temperatures can be obtained. For the SPSed and HPed samples, there is almost no plastic deformation at 500 ◦ C. When tested at 600 ◦ C, the TS values of SPSed and HPed samples are 536 and 425 MPa, which are higher than that of pure tungsten [32], while the TE values are 11.6% and 23.0%, respectively. Because the recrystallization occurred in the HP sintering process is more severe than in the SPS process as shown in Figures 6 and 7, the smaller grain size in the matrix can enhance the TS at comparative temperature, and the SPSed samples with the smaller grain size exhibit higher TS. As for the HPCRed samples, it exhibits almost no plastic deformation at 300 ◦ C, but the TS at 300 ◦ C is as high as 816 MPa. When tested at 400 ◦ C, the HPCRed samples undergo observable plastic deformation with a TE of 8.4% and a TS of 784 MPa. At 500 ◦ C, the TE further increase up to 17.5% and the TS reduces to 750 MPa. The fracture energy of the HPCRed samples is obviously larger than that of the SPSed and HPed samples, because the proportion of small angle grain boundaries is higher than others in both the fibres and matrix. The mechanical properties of the composites were studied preliminary, and other methods will be used to improve that.

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Figure 9. Tensile behavior of the Wf /W composites prepared with different methods between 300 ◦ C and 600 ◦ C.

The ductile fracture surfaces of the SPSed, HPed and HPCRed tensile tested specimens were shown in Figure 10. The SPSed (Figure 10a) and HPed samples (Figure 10b) show fibre cracking and necking at 600 ◦ C, which can be seen more clearly from the insets. In the HPCRed samples, fibre crack can be identified in Figure 10(c2) at 400 ◦ C and Figure 10(d2) at 500 ◦ C, whereas the matrix shows intergranular fracture behavior in both cases. This suggests that the enhanced tensile strength of the HPCRed samples over the SPSed and HPed samples can be mainly attributed to the elongated grains. However, a necking behavior instead of a plastic plateau after reaching the ultimate stress makes the HPCR method a bit defective, and whether a modified HPCR method is satisfied to obtain compact and tough Wf /W composites needs more research.

Figure 10. SEM micrographs of fracture surface in Wf /W composites tensile-tested at different temperatures: (a) SPS-600 ◦ C, (b) HP-600 ◦ C, (c) HPCR-400 ◦ C, and (d) HPCR-500 ◦ C. The insets on the top right corner of (a,b) are the enlarged views over a single fibre.

4. Conclusions Tungsten-fibre-net-reinforced tungsten composites (Wf /W) containing four layers of nets with enhanced fracture energy were synthesized by three different methods including spark plasma sintering (SPS), hot pressing (HP), and cold rolling after hot pressing (HPCR). The total thickness of the SPSed, HPed and HPCRed samples were 2.18, 3.52 and 2.08 mm, respectively, with corresponding fibre mass fraction of 17.4%, 10.5% and 10.5%, respectively. The microstructure, tensile property and fibre texture of these samples were investigated. The main results derived from such investigations can be concluded as follows: The relative density of all samples was above 95.10%, while the highest relative density of 99.80% is reached in the HPCRed samples. The hardness of the sintered fibres in all samples is around 537 HV0.2 which is smaller than the value of the original fibres (about 624 HV0.2 ). The proportion of grains with low misorientation angles 96%, Cl < 0.05%, Fe < 0.1%, K < 0.005%, Na < 0.005%, Ni < 0.005%, Pb < 0.005%, Zn < 0.005%) was supplied by Sigma-Aldrich, St. Louis, MO, USA. Pure Mg and Mg 2 vol % β-tricalcium phosphate (β-TCP) composite was synthesized using the powder metallurgy technique, incorporating hybrid microwave sintering [20]. The as-sintered billets were homogenized at 400 ◦ C for 1 h and were then hot extruded at 350 ◦ C to obtain cylindrical rods of 8 mm diameter at an extrusion ratio of 20.25:1. Samples that were cut from the rods were then characterized for physical and mechanical properties. 2.2. Material Characterization 2.2.1. Density Measurements Density measurements were performed on both monolithic and composite samples using the Archimedes principle. Four samples were cut from different parts of the extruded rods and were tested ten times for conformance. The samples were weighed separately in air and water using an A&D ER-182A electronic balance (Bradford, MA, USA) with an accuracy of 10−4 g. The theoretical density was calculated using the densities and weight percentages of the constituents by means of the rule of mixtures. From the experimental and theoretical densities, the porosity values of the samples were determined. 2.2.2. Microstructural Characterization Cylindrical samples were finely polished and etched according to the conventional techniques of metallography to obtain a clear distinction between the grain boundaries with the help of a LEICA-DM 2500M metallographic light microscope (Singapore). Four representative micrographs were analyzed for each composition in order to obtain accurate grain sizes. The OLYMPUS metallographic microscope (Singapore) and JEOL JSM-5800 LV Scanning Electron Microscope (SEM, Kyoto, Japan) was used for the microstructural characterization studies. X-ray diffraction studies were carried out on extruded samples in the direction along the axis of extrusion. The studies were performed using an automated Shimadzu LAB-XRD-6000 (Cu Kα; λ = 1.54056 Å, Kyoto, Japan) using a scan speed of 2◦ /min. The studies were conducted to identify the possible formation of any impurities/secondary phases. The XRD analysis was also conducted on the post-corroded samples to identify the corrosion products that were formed. 2.2.3. Damping and Elastic Modulus Damping characteristics and elastic modulus of the cylindrical samples (7 mm diameter and 60 mm length) were analyzed using the resonant frequency and the damping analyzer (RFDA) equipment from IMCE, Genk, Belgium. Recordings of the vibration signal were obtained in terms of amplitude vs. time. Damping capacity, loss rate, and elastic modulus values for both pure Mg and Mg (0.5, 1.0, and 1.5) vol % β-TCP composite sample were recorded.

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2.2.4. Mechanical Properties Microhardness tests were performed on the composite samples using Vickers microhardness tester Matsuzawa MXT 50 (Kyoto, Japan) with an indenter phase angle ~136◦ ; in conformance with ASTM standard E384-11-1 [21]. Fifteen readings were taken to arrive at an average representative value. Compression testing in the quasi-static mode was performed on cylindrical samples having 8 mm diameter and 8 mm length, utilizing a fully automated servo-hydraulic mechanical testing machine (Model-MTS 810; in conformance with ASTM test method E9-09, Eden Prairie, MN, USA) at a strain rate 8.33 × 10−5 s−1 [22]. Four specimens each for both of the compositions were tested to ensure reproducibility. Fracture surface analysis of the samples failed under compression was done using SEM. 2.2.5. Immersion Studies Cylindrical samples of (5 mm diameter and 5 mm length) were immersed for 96 h in Hanks balanced salt solution (HBSS) procured from Lonza Chemicals Pte Ltd. Singapore. The setup was immersed in a water bath that was maintained at 37 ◦ C to simulate the temperature of the human body. The sample dimensions of 5 mm diameter and 5 mm length was used. Solution to sample ratio was maintained at 20 mL:1 cm2 . The solution was changed every 24 h. Weight loss and pH measurements were measured after every 24 h. A solution containing 20 g CrO3 and 1.9 g AgNO3 dissolved in 100 mL of de-ionized water was used for removing the corrosion products. The samples post-corrosion were observed under the SEM in order to gain further information about the nature of corrosion products that were formed. 3. Results and Discussion 3.1. Density and Porosity Table 1 shows the density and porosity levels of pure Mg and Mg-β-TCP composites. The experimental density of Pure Mg slightly increased with the incorporation of β-TCP, and Mg-1.5 TCP composite exhibited an experimental density value of 1.7449 g·cm−3 . The slight increase (0.2%) in the density can be attributed to the fact that there is a density difference between the matrix (1.74 g·cm−3 ) and reinforcement (3.14 g·cm−3 ). Porosity levels marginally increased with the addition of the β-TCP and the highest porosity value of ~0.28% was observed for the Mg-1.5 TCP composite. The observed porosity is less than 1% porosity, which is an advantage when compared to conventional sintering processes that can achieve only up to ~95% densification [23]. Microstructural examination of the extruded rod revealed the absence of blowholes, defects and a superior surface finish indicated the suitability of the powder metallurgy technique to generate near dense composites [3]. Table 1. Density, porosity, grain size and microhardness measurements of pure Mg and Mg-β-tricalcium phosphate (β-TCP) composite. Material

Theoretical Density (g cm−3 )

Experimental Density (g cm−3 )

Porosity (%)

Grain Size (μm)

Hardness (Hv)

Pure Mg Mg-0.5 TCP Mg-1.0 TCP Mg-1.5 TCP

1.74 1.7412 1.7424 1.7449

1.7363 ± 0.002 1.7371 ± 0.0147 1.7381 ± 0.0067 1.7387 ± 0.0048

0.21 0.23 0.24 0.28

34 ± 2 18 ± 2 (↓47%) 13 ± 1 (↓61%) 10 ± 1 (↓70%)

46 ± 3 52 ± 2 (↑13.04%) 54 ± 3 (↑17.39%) 54 ± 1 (↑17.39%)

3.2. Microstructural Characterisation Table 1 shows the average grain size values of pure Mg and Mg (0.5, 1.0, and 1.5) vol % β-TCP composites. The grain size of pure Mg in as-extruded form was observed to be ~34 μm. The addition of 0.5 vol % β-TCP particles resulted in superior grain refinement of up to ~18 μm, which is ~47% finer than that of pure Mg. Increased addition of 1.0 and 1.5 vol % β-TCP particles resulted in a further

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refinement in the grain size of up to ~13 and ~10 μm. Near equiaxed grain morphology was observed with the addition of β-TCP particles, as observed in Figure 1. This superior grain refinement can be attributed to mainly two aspects namely (a) Particle stimulated nucleation phenomenon that promotes the nucleation of grains, hence restricting the grain growth; (b) the ability of the β-TCP particles to pin the grain boundaries resulting in finer microstructure [24]. As the size of the reinforcement is predominantly in micron length scale, simulated dynamic recrystallization phenomenon can be expected during the extrusion process [25]. Distribution of the β-TCP particle in the Mg matrix is shown in Figure 2. The efficient extrusion process has managed to break down the large β-TCP particles and clusters, leading to a reasonable distribution pattern. Hence, hot extrusion can be considered as a suitable secondary process to promote the near uniform distribution of reinforcement and simultaneously reducing the spatial heterogeneity of the mechanical properties of the Mg-based composites [16]. The superior grain refinement also aids in the strengthening of the composites by means of Hall-Petch mechanism activation. High wettability of β-TCP particles with the Mg matrix leads to the easy densification under sintering, and hence showing a superior interfacial integrity between the particle and the matrix. The near-uniform distribution of β-TCP throughout the Mg matrix can also be attributed to the suitable primary and secondary processing parameters that are optimized for the primary processing of Mg-β-TCP composites. Energy Dispersive Spectroscopy (EDS) analysis of Mg-1.0 TCP composite is also shown in Figure 2. The EDS spectra are studied at the reinforcement (A) and matrix (B) location. The analysis of the matrix reveals predominantly Mg phases with traces of O due to surface oxidation during the processing of the material. The β-TCP particles have settled at the grain boundaries of the composite, hence confirming the grain boundary pinning mechanism that is responsible for grain refinement as quantitatively confirmed by the predominant Ca, O, P peaks in the spectrum. The EDS also confirms the absence of any sign of impurities or secondary phases in the composite.

Figure 1. Optical micrography images of Mg-β-TCP composites: (a) Mg-0.5 TCP; (b) Mg-1.0 TCP; and, (c) Mg-1.5 TCP.

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Figure 2. (a) Grain boundary pinning mechanism of Mg-1.0 TCP composite; (b) β-TCP particle distribution within the Mg matrix in Mg-1.5 TCP composite. Energy Dispersive Spectroscopy (EDS) analysis of the Mg-1.0 TCP composite at the matrix and reinforcement location.

The X-ray diffraction analysis of the developed composites was performed along the extruded direction and shown in Figure 3. The X-ray diffraction peaks of pure Mg and Mg-β-TCP composites reveal mainly Mg peaks. The reinforcement peaks are not visible, as the amount of reinforcement in the Mg matrix is low (